Process for forming low defect density, ideal oxygen precipitating silicon

ABSTRACT

The present invention is directed to a process for producing a silicon wafer which, during the heat treatment cycles of essentially any arbitrary electronic device manufacturing process, may form an ideal, non-uniform depth distribution of oxygen precipitates and may additionally contain an axially symmetric region which is substantially free of agglomerated intrinsic point defects. The process either comprises exposing the wafer&#39;s front and back surfaces to different atmospheres, or thermally annealing two wafers in a face-to-face arrangement.

REFERENCE TO RELATED APPLICATION

This is a continuation of application Ser. No. 11/058,996, filed on Feb.16, 2005; which is a continuation of Ser. No. 10/373,899, filed on Feb.25, 2003, now U.S. Pat. No. 6,896,728; which is a division ofapplication Ser. No. 09/705,092, filed Nov. 2, 2000, now U.S. Pat. No.6,555,194; which is a continuation of application Ser. No. 09/057,800,filed Apr. 9, 1998, now U.S. Pat. No. 6,190,631; which claims thebenefit of U.S. Provisional Application No. 60/041,845, filed Apr. 9,1997, and U.S. Provisional Application No. 60/062,316, filed Oct. 17,1997; all of which are incorporated in their entirety herein byreference.

BACKGROUND OF THE INVENTION

The present invention generally relates to the preparation ofsemiconductor material substrates, especially silicon wafers, which areused in the manufacture of electronic components. More particularly, thepresent invention relates to a silicon wafer, and a process for thepreparation thereof, which, during the heat treatment cycles ofessentially any arbitrary electronic device manufacturing process, mayform an ideal, non-uniform depth distribution of oxygen precipitates andmay additionally contain an axially symmetric region which issubstantially free of agglomerated intrinsic point defects.

Single crystal silicon, which is the starting material for mostprocesses for the fabrication of semiconductor electronic components, iscommonly prepared with the so-called Czochralski process wherein asingle seed crystal is immersed into molten silicon and then grown byslow extraction. As molten silicon is contained in a quartz crucible, itis contaminated with various impurities, among which is mainly oxygen.At the temperature of the silicon molten mass, oxygen comes into thecrystal lattice until it reaches a concentration determined by thesolubility of oxygen in silicon at the temperature of the molten massand by the actual segregation coefficient of oxygen in solidifiedsilicon. Such concentrations are greater than the solubility of oxygenin solid silicon at the temperatures typical for the processes for thefabrication of electronic devices. As the crystal grows from the moltenmass and cools, therefore, the solubility of oxygen in it decreasesrapidly, whereby in the resulting slices or wafers, oxygen is present insupersaturated concentrations.

Thermal treatment cycles which are typically employed in the fabricationof electronic devices can cause the precipitation of oxygen in siliconwafers which are supersaturated in oxygen. Depending upon their locationin the wafer, the precipitates can be harmful or beneficial. Oxygenprecipitates located in the active device region of the wafer can impairthe operation of the device. Oxygen precipitates located in the bulk ofthe wafer, however, are capable of trapping undesired metal impuritiesthat may come into contact with the wafer. The use of oxygenprecipitates located in the bulk of the wafer to trap metals is commonlyreferred to as internal or intrinsic gettering (“IG”).

Historically, electronic device fabrication processes included a seriesof steps which were designed to produce silicon having a zone or regionnear the surface of the wafer which is free of oxygen precipitates(commonly referred to as a “denuded zone” or a “precipitate free zone”)with the balance of the wafer, i.e., the wafer bulk, containing asufficient number of oxygen precipitates for IG purposes. Denuded zonescan be formed, for example, in a high-low-high thermal sequence such as(a) oxygen out-diffusion heat treatment at a high temperature (>1100°C.) in an inert ambient for a period of at least about 4 hours, (b)oxygen precipitate nuclei formation at a low temperature (600-750° C.),and (c) growth of oxygen (SiO₂) precipitates at a high temperature(1000-1150° C.). See, e.g., F. Shimura, Semiconductor Silicon CrystalTechnology, Academic Press, Inc., San Diego Calif. (1989) at pages361-367 and the references cited therein.

More recently, however, advanced electronic device manufacturingprocesses such as DRAM manufacturing processes have begun to minimizethe use of high temperature process steps. Although some of theseprocesses retain enough of the high temperature process steps to producea denuded zone and sufficient density of bulk precipitates, thetolerances on the material are too tight to render it a commerciallyviable product. Other current highly advanced electronic devicemanufacturing processes contain no out-diffusion steps at all. Becauseof the problems associated with oxygen precipitates in the active deviceregion, therefore, these electronic device fabricators must use siliconwafers which are incapable of forming oxygen precipitates anywhere inthe wafer under their process conditions. As a result, all IG potentialis lost.

SUMMARY OF THE INVENTION

Among the objects of the invention, therefore, is the provision of aprocess for heat-treating a Cz, single crystal silicon wafer toinfluence the precipitation behavior of oxygen in the wafer in asubsequent thermal processing step, the wafer comprising two major,generally parallel surfaces, one of which is a front surface of thewafer and the other of which is a back surface of the wafer, a centralplane between the front and back surfaces, a circumferential edgejoining the front and back surfaces, a central axis generallyperpendicular to the front and back surfaces, a radius extending fromthe central axis to the circumferential edge, a nominal diameter of 150mm, 200 mm, or greater than 200 mm, a front surface layer whichcomprises the region of the wafer between the front surface and adistance, D₁, of at least about 10 micrometers measured from the frontsurface and toward the central plane, a back surface layer whichcomprises a second region of the wafer between the back surface and adistance D₂, of at least about 10 micrometers as measured from the backsurface and toward the central plane, and a bulk layer which comprises athird region of the wafer between the front surface layer and the backsurface layer. The wafer further comprises an axially symmetric regionwhich has radial width of at least about 30% the length of the radius ofthe wafer and has no detectible agglomerated intrinsic point defects ata detection limit of 10³ defects/cm³. The process comprises the steps ofsubjecting the wafer to a heat-treatment in excess of about 1175° C. ina non-oxidizing atmosphere to form crystal lattice vacancies in thefront surface and bulk layers, and controlling the cooling rate of theheat-treated wafer to produce a wafer having a vacancy concentrationprofile in which a maximum concentration is located between the centralplane and the back surface layer, the vacancy concentration generallyincreasing from the front surface to the region of maximum concentrationand generally decreasing from the region of maximum concentration to theback surface with the difference in the concentration of vacancies inthe front surface layer, the back surface layer, and the bulk layerbeing such that a thermal treatment at a temperature in excess of 750°C. is capable of forming in the wafer a denuded zone in the frontsurface layer and in the back surface layer and oxygen clusters orprecipitates in the bulk zone, with the concentration of the oxygenclusters or precipitates in the bulk layer being primarily dependantupon the concentration of vacancies.

The invention is further directed to a process for producing a singlecrystal silicon wafer having two major, generally parallel surfaces, oneof which is a front surface of the wafer and the other of which is aback surface of the wafer, a central plane between the front and backsurfaces, a circumferential edge joining the front and back surfaces, acentral axis generally perpendicular to the front and back surfaces, aradius extending from the central axis to the circumferential edge, anominal diameter of 150 mm, 200 mm, or greater than 200 mm, a frontsurface layer which comprises the region of the wafer between the frontsurface and a distance, D₁, of at least about 10 micrometers measuredfrom the front surface and toward the central plane, a back surfacelayer which comprises a second region of the wafer between the backsurface and a distance D₂, of at least about 10 micrometers as measuredfrom the back surface and toward the central plane, and a bulk layerwhich comprises a third region of the wafer between the front surfacelayer and the back surface layer. The wafer has a non-uniformdistribution of crystal lattice vacancies in which a maximumconcentration is located between the central plane and the back surfacelayer, the vacancy concentration generally increasing from the frontsurface to the region of maximum concentration and generally decreasingfrom the region of maximum concentration to the back surface, the waferfurther comprising a first axially symmetric region, which has nodetectible agglomerated intrinsic point defects at a detection limit of10³ defects/cm³. The process comprising the steps of growing a singlecrystal silicon ingot in which the ingot comprises a central axis, aseed-cone, an end-cone and a constant diameter portion between theseed-cone and the end-cone having a circumferential edge, a radiusextending from the central axis to the circumferential edge, and anominal diameter of 150 mm, 200 mm, or greater than 200 mm; controlling(i) a growth velocity, v, (ii) an average axial temperature gradient,G₀, during the growth of the constant diameter portion of the crystalover the temperature range from solidification to a temperature of noless than about 1325° C., and (iii) the cooling rate of the crystal fromthe solidification temperature to about 1,050° C. to cause the formationof an axially symmetrical segment which has no detectible agglomeratedintrinsic point defects at a detection limit of 10³ defects/cm³, whereinthe axially symmetric region extends inwardly from the circumferentialedge of the ingot, has a width as measured from the circumferential edgeradially toward the central axis of the ingot which is at least about30% the length of the radius of the ingot, and has a length as measuredalong the central axis of at least about 20% the length of the constantdiameter portion of the ingot; slicing a single crystal silicon waferfrom the constant diameter portion of the ingot; subjecting the wafer toa heat-treatment in excess of about 1175° C. in a non-oxidizingatmosphere to form crystal lattice vacancies in the front surface andbulk layers; and, controlling the cooling rate of the heat-treated waferto produce a wafer having a vacancy concentration profile in which amaximum concentration is located between the central plane and the backsurface layer, the vacancy concentration generally increasing from thefront surface to the region of maximum concentration and generallydecreasing from the region of maximum concentration to the back surfaceand the difference in the concentration of vacancies in the frontsurface layer, the back surface layer, and the bulk layer being suchthat a thermal treatment at a temperature in excess of 750° C. iscapable of forming in the wafer a denuded zone in the front surfacelayer and in the back surface layer and oxygen clusters or precipitatesin the bulk zone, with the concentration of the oxygen clusters orprecipitates in the bulk layer being primarily dependant upon theconcentration of vacancies.

The present invention is still further directed to a process forheat-treating a Cz, single crystal silicon wafer, to influence theprecipitation behavior of oxygen in the wafer in a subsequent thermalprocessing step, the wafer comprising two major, generally parallelsurfaces, one of which is a front surface of the wafer and the other ofwhich is a back surface of the wafer, a central plane between the frontand back surfaces, a circumferential edge joining the front and backsurfaces, a central axis generally perpendicular to the front and backsurfaces, a radius extending from the central axis to thecircumferential edge, a nominal diameter of 150 mm, 200 mm, or greaterthan 200 mm, a front surface layer which comprises the region of thewafer between the front surface and a distance, D₁, of at least about 10micrometers measured from the front surface and toward the centralplane, a back surface layer which comprises a second region of the waferbetween the back surface and a distance D₂, of at least about 10micrometers as measured from the back surface and toward the centralplane, and a bulk layer which comprises a third region of the waferbetween the front surface layer and the back surface layer. The waferfurther comprises a first axially symmetric region in which vacanciesare the predominant intrinsic point defect and which has no detectibleagglomerated intrinsic point defects at a detection limit of 10³defects/cm³, wherein the first axially symmetric region comprises thecentral axis or has a width of at least about 15 mm. The processcomprises the steps of subjecting the wafer to a heat-treatment inexcess of about 1175° C. in a non-oxidizing atmosphere to form crystallattice vacancies in the surface layer and the bulk layer; and,controlling the cooling rate of the heat-treated wafer to produce awafer having a non-uniform concentration of vacancies in which a maximumconcentration is located between the central plane and the back surfacelayer, the vacancy concentration generally increasing from the frontsurface to the region of maximum concentration and generally decreasingfrom the region of maximum concentration to the back surface such that,upon subjecting the wafer to an oxygen precipitation heat treatment, adenuded zone is formed in the surface layer and oxygen clusters orprecipitates are formed in the bulk layer with the concentration of theoxygen clusters or precipitates in the bulk layer being primarilydependant upon the concentration of vacancies.

Other objects and features of this invention will be in part apparentand in part pointed out hereinafter.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic depiction of the process of the present invention.

FIG. 2 is a photograph of a cross-section of a wafer (sample 4-7) whichwas processed as described in Example 1.

FIG. 3 is a photograph of a cross-section of a wafer (sample 4-8) whichwas subjected to the series of steps described in Example 1.

FIG. 4 is a photograph of a cross-section of a wafer (sample 3-14) whichwas subjected to the series of steps described in Example 1.

FIG. 5 is a graph of the log of platinum concentration (atoms/cm³)versus depth from the surface of a wafer (sample 4-7) which wassubjected to the series of steps set forth in Example 1.

FIG. 6 is a photograph of a cross-section of a wafer (sample 3-4) whichwas subjected to the series of steps set forth in Example 2.

FIG. 7 is a photograph of a cross-section of a wafer (sample 3-5) whichwas subjected to the series of steps set forth in Example 2.

FIG. 8 is a photograph of a cross-section of a wafer (sample 3-6) whichwas subjected to the series of steps set forth in Example 2.

FIG. 9 is a photograph of a cross-section of a wafer (sample 1-8) whichwas subjected to the series of steps set forth in Example 3.

FIG. 10 is logarithmic graph of the number density of bulk microdefects(BMD) versus the partial pressure of oxygen present in the atmosphereduring rapid thermal annealing of single crystal silicon wafers inaccordance with the present invention, as described in Example 4.

FIG. 11 is a graph which shows an example of how the initialconcentration of self-interstitials, [I], and vacancies, [V], changeswith an increase in the value of the ratio v/G₀, where v is the growthrate and G₀ is the average axial temperature gradient.

FIG. 12 is a graph which shows an example of how ΔG_(I), the change infree energy required for the formation of agglomerated interstitialdefects, increases as the temperature, T, decreases, for a given initialconcentration of self-interstitials, [I].

FIG. 13 is a graph which shows an example of how the initialconcentration of self-interstitials, [I], and vacancies, [V], can changealong the radius of an ingot or wafer, as the value of the ratio v/G₀decreases, due to an increase in the value of G₀. Note that at the V/Iboundary a transition occurs from vacancy dominated material toself-interstitial dominated material.

FIG. 14 is a top plan view of a single crystal silicon ingot or wafershowing regions of vacancy, V, and self-interstitial, I, dominatedmaterials respectively, as well as the V/I boundary that exists betweenthem.

FIG. 15 is a longitudinal, cross-sectional view of a single crystalsilicon ingot showing, in detail, an axially symmetric region of aconstant diameter portion of the ingot.

FIG. 16 is an image produced by a scan of the minority carrier lifetimeof an axial cut of the ingot following a series of oxygen precipitationheat treatments, showing in detail a generally cylindrical region ofvacancy dominated material, a generally annular shaped axially symmetricregion of self-interstitial dominated material, the V/I boundary presentbetween them, and a region of agglomerated interstitial defects.

FIG. 17 is a graph of pull rate (i.e. seed lift) as a function ofcrystal length, showing how the pull rate is decreased linearly over aportion of the length of the crystal.

FIG. 18 is an image produced by a scan of the minority carrier lifetimeof an axial cut of the ingot following a series of oxygen precipitationheat treatments, as described in Example 5.

FIG. 19 is a graph of pull rate as a function of crystal length for eachof four single crystal silicon ingots, labeled 1-4 respectively, whichare used to yield a curve, labeled v*(Z), as described in Example 5.

FIG. 20 is a graph of the average axial temperature gradient at themelt/solid interface, G₀, as a function of radial position, for twodifferent cases as described in Example 6.

FIG. 21 is a graph of the initial concentration of vacancies, [V], orself-interstitials, [I], as a function of radial position, for twodifferent cases as described Example 6.

FIG. 22 is a graph of temperature as a function of axial position,showing the axial temperature profile in ingots for two different casesas described in Example 7.

FIG. 23 is a graph of the self-interstitial concentrations resultingfrom the two cooling conditions illustrated in FIG. 22 and as more fullydescribed in Example 7.

FIG. 24 is an image produced by a scan of the minority carrier lifetimeof an axial cut of an entire ingot following a series of oxygenprecipitation heat treatments, as described in Example 8.

FIG. 25 is a graph illustrating the position of the V/I boundary as afunction of the length of the single crystal silicon ingot, as describedin Example 9.

FIG. 26 a is an image produced by a scan of the minority carrierlifetime of an axial cut of a segment of an ingot, ranging from about100 mm to about 250 mm from the shoulder of the ingot, following aseries of oxygen precipitation heat treatments, as described in Example10.

FIG. 26 b is an image produced by a scan of the minority carrierlifetime of an axial cut of a segment of an ingot, ranging from about250 mm to about 400 mm from the shoulder of the ingot, following aseries of oxygen precipitation heat treatments, as described in Example10.

FIG. 27 is a graph of the axial temperature gradient, G₀, at variousaxial positions for an ingot, as described in Example 11.

FIG. 28 is a graph of the radial variations in the average axialtemperature gradient, G₀, at various for an ingot, as described inExample 11.

FIG. 29 is a graph illustrating the relationship between the width ofthe axially symmetric region and the cooling rate, as described inExample 11.

FIG. 30 is a photograph of an axial cut of a segment of an ingot,ranging from about 235 mm to about 350 mm from the shoulder of theingot, following copper decoration and a defect-delineating etch,described in Example 11.

FIG. 31 is a photograph of an axial cut of a segment of an ingot,ranging from about 305 mm to about 460 mm from the shoulder of theingot, following copper decoration and a defect-delineating etch,described in Example 11.

FIG. 32 is a photograph of an axial cut of a segment of an ingot,ranging from about 140 mm to about 275 mm from the shoulder of theingot, following copper decoration and a defect-delineating etch,described in Example 11.

FIG. 33 is a photograph of an axial cut of a segment of an ingot,ranging from about 600 mm to about 730 mm from the shoulder of theingot, following copper decoration and a defect-delineating etch,described in Example 11.

FIG. 34 is a graph illustrating the radial variations in the averageaxial temperature gradient, G₀(r), which may occur in hot zones ofvarious configurations.

FIG. 35 is a graph illustrating the axial temperature profile for aningot in four different hot zone configurations.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

In accordance with the present invention, an ideal precipitating waferhas been discovered which, during essentially any electronic devicemanufacturing process, will form a denuded zone of sufficient depth anda wafer bulk containing a sufficient density of oxygen precipitates forIG purposes. Advantageously, this ideal precipitating wafer may beprepared in a matter of minutes using tools which are in common use inthe semiconductor silicon manufacturing industry. This process creates a“template” in the silicon which determines or “prints” the manner inwhich oxygen will precipitate during the electronic device manufacturingprocess.

The starting material for the ideal precipitating wafer of the presentinvention is a single crystal silicon wafer which has been sliced from asingle crystal ingot grown in accordance with conventional Czochralskicrystal growing methods. Such methods, as well as standard siliconslicing, lapping, etching, and polishing techniques are disclosed, forexample, in F. Shimura, Semiconductor Silicon Crystal Technology,Academic Press, 1989, and Silicon Chemical Etching, (J. Grabmaier ed.)Springer-Verlag, N.Y., 1982 (incorporated herein by reference).

Czochralski-grown silicon typically has an oxygen concentration withinthe range of about 5×10¹⁷ to about 9×10¹⁷ atoms/cm³ (ASTM standardF-121-83). Because the oxygen precipitation behavior of the waferbecomes essentially decoupled from the oxygen concentration in the idealprecipitating wafer, the starting wafer may have an oxygen concentrationfalling anywhere within or even outside the range attainable by theCzochralski process.

Depending upon the cooling rate of the single crystal silicon ingot fromthe temperature of the melting point of silicon (about 1410° C.) throughthe range of about 750° C. to about 350° C., oxygen precipitatenucleation centers may form in the single crystal silicon ingot fromwhich the wafer is sliced. The presence or absence of these nucleationcenters in the starting material is not critical to the presentinvention provided, however, these centers are capable of beingdissolved by heat-treating the silicon at temperatures not in excess ofabout 1300° C. Certain heat-treatments, such as annealing the silicon ata temperature of about 800° C. for about four hours, can stabilize thesecenters such that they are incapable of being dissolved at temperaturesnot in excess of about 1150° C. The detection limit for oxygenprecipitates is currently about 5×10⁶ precipitates/cm³. The presence (ordensity) of oxygen precipitation nucleation centers cannot be directlymeasured using presently available techniques. Various techniques may beused, however, to indirectly detect their presence. As previouslydiscussed, preexisting oxygen precipitate nucleation centers in thesilicon can be stabilized and precipitates can be grown at these sitesby subjecting the silicon to an oxygen precipitation heat treatment.Thus, the presence of these nucleation centers can indirectly bemeasured after an oxygen precipitation heat treatment, e.g., annealingthe wafer at a temperature of 800° C. for four hours and then at atemperature of 1000° C. for sixteen hours.

Substitutional carbon, when present as an impurity in single crystalsilicon, has the ability to catalyze the formation of oxygen precipitatenucleation centers. For this and other reasons, therefore, it ispreferred that the single crystal silicon starting material have a lowconcentration of carbon. That is, the single crystal silicon should havea concentration of carbon which is less than about 5×10¹⁶ atoms/cm³,preferably which is less than 1×10¹⁶ atoms/cm³, and more preferably lessthan 5×10¹⁵ atoms/cm³.

Referring now to FIG. 1, the starting material for the idealprecipitating wafer of the present invention, single crystal siliconwafer 1, has a front surface 3, a back surface 5, and an imaginarycentral plane 7 between the front and back surfaces. The terms “front”and “back” in this context are used to distinguish the two major,generally planar surfaces of the wafer; the front surface of the waferas that term is used herein is not necessarily the surface onto which anelectronic device will subsequently be fabricated nor is the backsurface of the wafer as that term is used herein necessarily the majorsurface of the wafer which is opposite the surface onto which theelectronic device is fabricated. In addition, because silicon waferstypically have some total thickness variation (TTV), warp and bow, themidpoint between every point on the front surface and every point on theback surface may not precisely fall within a plane; as a practicalmatter, however, the TTV, warp and bow are typically so slight that to aclose approximation the midpoints can be said to fall within animaginary central plane which is approximately equidistant between thefront and back surfaces.

In a first embodiment of the process of the present invention wafer 1 isheat-treated in an oxygen-containing atmosphere in step S₁ to grow asuperficial oxide layer 9 which envelopes wafer 1. In general, the oxidelayer will have a thickness which is greater than the native oxide layerwhich forms upon silicon (about 15 Ångstroms); preferably, the oxidelayer has a thickness of at least about 20 Ångstroms and, in someembodiments, at least about 25 Ångstroms or even at least about 30Ångstroms. Experimental evidence obtained to-date, however, suggeststhat oxide layers having a thickness greater than about 30 Ångstroms,while not interfering with the desired effect, provide little or noadditional benefit.

In step S₂, the wafer is subjected to a heat-treatment step in which thewafers are heated to an elevated temperature to form and therebyincrease the number density of crystal lattice vacancies 13 in wafer 1.Preferably, this heat-treatment step is carried out in a rapid thermalannealer in which the wafers are rapidly heated to a target temperatureand annealed at that temperature for a relatively short period of time.In general, the wafer is subjected to a temperature in excess of 1150°C., preferably at least 1175° C., more preferably at least about 1200°C., and most preferably between about 1200° C. and 1275° C.

In the first embodiment of the present invention, the rapid thermalannealing step is carried out in the presence of a nitriding atmosphere,that is, an atmosphere containing nitrogen gas (N₂) or anitrogen-containing compound gas such as ammonia which is capable ofnitriding an exposed silicon surface. The atmosphere may thus consistentirely of nitrogen or nitrogen compound gas, or it may additionallycomprise a non-nitriding gas such as argon. An increase in vacancyconcentration throughout the wafer is achieved nearly, if notimmediately, upon achieving the annealing temperature. The wafer willgenerally be maintained at this temperature for at least one second,typically for at least several seconds (e.g., at least 3), preferablyfor several tens of seconds (e.g., 20, 30, 40, or 50 seconds) and,depending upon the desired characteristics of the wafer, for a periodwhich may range up to about 60 seconds (which is near the limit forcommercially available rapid thermal annealers). The resulting waferwill have a relatively uniform vacancy concentration (number density)profile in the wafer.

Based upon experimental evidence obtained to-date, the atmosphere inwhich the rapid thermal annealing step is carried out preferably has nomore than a relatively small partial pressure of oxygen, water vapor andother oxidizing gases; that is, the atmosphere has a total absence ofoxidizing gases or a partial pressure of such gases which isinsufficient to inject sufficient quantities of siliconself-interstitial atoms which suppress the build-up of vacancyconcentrations. While the lower limit of oxidizing gas concentration hasnot been precisely determined, it has been demonstrated that for partialpressures of oxygen of 0.01 atmospheres (atm.), or 10,000 parts permillion atomic (ppma), no increase in vacancy concentration and noeffect is observed. Thus, it is preferred that the atmosphere have apartial pressure of oxygen and other oxidizing gases of less than 0.01atm. (10,000 ppma); more preferably the partial pressure of these gasesin the atmosphere is no more than about 0.005 atm. (5,000 ppma), morepreferably no more than about 0.002 atm. (2,000 ppma), and mostpreferably no more than about 0.001 atm. (1,000 ppma).

In addition to causing the formation of crystal lattice vacancies, therapid thermal annealing step causes the dissolution of any unstabilizedoxygen precipitate nucleation centers which are present in the siliconstarting material. These nucleation centers may be formed, for example,during the growth of the single crystal silicon ingot from which thewafer was sliced, or as a consequence of some other event in theprevious thermal history of the wafer or of the ingot from which thewafer is sliced. Thus, the presence or absence of these nucleationcenters in the starting material is not critical, provided these centersare capable of being dissolved during the rapid thermal annealing step.

The rapid thermal anneal may be carried out in any of a number ofcommercially available rapid thermal annealing (“RTA”) furnaces in whichwafers are individually heated by banks of high power lamps. RTAfurnaces are capable of rapidly heating a silicon wafer, e.g., they arecapable of heating a wafer from room temperature to 1200° C. in a fewseconds. One such commercially available RTA furnace is the model 610furnace available from AG Associates (Mountain View, Calif.).

Intrinsic point defects (vacancies and silicon self-interstitials) arecapable of diffusing through single crystal silicon with the rate ofdiffusion being temperature dependant. The concentration profile ofintrinsic point defects, therefore, is a function of the diffusivity ofthe intrinsic point defects and the recombination rate as a function oftemperature. For example, the intrinsic point defects are relativelymobile at temperatures in the vicinity of the temperature at which thewafer is annealed in the rapid thermal annealing step whereas they areessentially immobile for any commercially practical time period attemperatures of as much as 700° C. Experimental evidence obtainedto-date suggests that the effective diffusion rate of vacancies slowsconsiderably at temperatures less than about 700° C. and perhaps asgreat as 800° C., 900° C., or even 1,000° C., the vacancies can beconsidered to be immobile for any commercially practical time period.

Upon completion of step S₂, the wafer is rapidly cooled in step S₃through the range of temperatures at which crystal lattice vacancies arerelatively mobile in the single crystal silicon. As the temperature ofthe wafer is decreased through this range of temperatures, the vacanciesdiffuse to the oxide layer 9 and become annihilated, thus leading to achange in the vacancy concentration profile with the extent of changedepending upon the length of time the wafer is maintained at atemperature within this range. If the wafer were held at thistemperature within this range for an infinite period of time, thevacancy concentration would once again become substantially uniformthroughout wafer bulk 11 with the concentration being an equilibriumvalue which is substantially less than the concentration of crystallattice vacancies immediately upon completion of the heat treatmentstep. By rapidly cooling the wafer, however, a non-uniform distributionof crystal lattice vacancies can be achieved with the maximum vacancyconcentration being at or near central plane 7 and the vacancyconcentration decreasing in the direction of the front surface 3 andback surface 5 of the wafer. In general, the average cooling rate withinthis range of temperatures is at least about 5° C. per second andpreferably at least about 20° C. per second. Depending upon the desireddepth of the denuded zone, the average cooling rate may preferably be atleast about 50° C. per second, still more preferably at least about 100°C. per second, with cooling rates in the range of about 100° C. to about200° C. per second being presently preferred for some applications. Oncethe wafer is cooled to a temperature outside the range of temperaturesat which crystal lattice vacancies are relatively mobile in the singlecrystal silicon, the cooling rate does not appear to significantlyinfluence the precipitating characteristics of the wafer and thus, doesnot appear to be narrowly critical. Conveniently, the cooling step maybe carried out in the same atmosphere in which the heating step iscarried out.

In step S₄, the wafer is subjected to an oxygen precipitationheat-treatment. For example, the wafer may be annealed at a temperatureof 800° C. for four hours and then at a temperature of 1000° C. forsixteen hours. Alternatively and preferably, the wafer is loaded into afurnace which is at a temperature of about 800° C. as the first step ofan electronic device manufacturing process. When loaded into a furnaceat this temperature, the previously rapidly thermal annealed wafer willhave separate zones which behave differently with respect to oxygenprecipitation. In the high vacancy regions (the wafer bulk), oxygenclusters rapidly as the wafer enters the furnace. By the time theloading temperature is reached, the clustering process is finished and adistribution of clusters is reached which depends only upon the initialconcentration of vacancies. In the low vacancy regions (near the wafersurfaces), the wafer behaves like a normal wafer which lackspre-existing oxygen precipitate nucleation centers; that is, oxygenclustering is not observed. As the temperature is increased above 800°C. or if the temperature remains constant, the clusters in the vacancyrich zone grow into precipitates and are thereby consumed, whereas inthe vacancy lean zone nothing happens. By dividing the wafer intovarious zones of vacancy concentration, a template is effectivelycreated through which is written an oxygen precipitate pattern which isfixed the moment the wafer is loaded into the furnace.

As illustrated in FIG. 1, the resulting depth distribution of oxygenprecipitates in the wafer is characterized by clear regions of oxygenprecipitate-free material (denuded zones) 15 and 15′ extending from thefront surface 3 and back surface 5 to a depth t, t′, respectively.Between the oxygen precipitate-free regions, 15 and 15′, there is aregion 17 which contains a substantially uniform density of oxygenprecipitates.

The concentration of oxygen precipitates in region 17 is primarily afunction of the heating step and secondarily a function of the coolingrate. In general, the concentration of oxygen precipitates increaseswith increasing temperature and increasing annealing times in theheating step, with precipitate densities in the range of about 1×10⁷ toabout 5×10¹⁰ precipitates/cm³ being routinely obtained.

The depth t, t′ from the front and back surfaces, respectively, ofoxygen precipitate-free material (denuded zones) 15 and 15′ is primarilya function of the cooling rate through the temperature range at whichcrystal lattice vacancies are relatively mobile in silicon. In general,the depth t, t′ increases with decreasing cooling rates, with denudedzone depths of at least about 10, 20, 30, 40, 50, 70 or even 100micrometers being attainable. Significantly, the depth of the denudedzone is essentially independent of the details of the electronic devicemanufacturing process and, in addition, does not depend upon theout-diffusion of oxygen as is conventionally practiced.

While the rapid thermal treatments employed in this process of thepresent invention may result in the out-diffusion of a small amount ofoxygen from the surface of the front and back surfaces of the wafer, theamount of out-diffusion is significantly less than what is observed inconventional processes for the formation of denuded zones. As a result,the ideal precipitating wafers of the present invention have asubstantially uniform interstitial oxygen concentration as a function ofdistance from the silicon surface. For example, prior to the oxygenprecipitation heat-treatment, the wafer will have a substantiallyuniform concentration of interstitial oxygen from the center of thewafer to regions of the wafer which are within about 15 microns of thesilicon surface, more preferably from the center of the silicon toregions of the wafer which are within about 10 microns of the siliconsurface, even more preferably from the center of the silicon to regionsof the wafer which are within about 5 microns of the silicon surface,and most preferably from the center of the silicon to regions of thewafer which are within about 3 microns of the silicon surface. In thiscontext, a substantially uniform oxygen concentration shall mean avariance in the oxygen concentration of no more than about 50%,preferably no more than about 20%, and most preferably no more thanabout 10%.

Typically, oxygen precipitation heat-treatments do not result in asubstantial amount of oxygen outdiffusion from the heat-treated wafer.As a result, the concentration of interstitial oxygen in the denudedzone at distances more than several microns from the wafer surface willnot significantly change as a consequence of the precipitationheat-treatment. For example, if the denuded zone of the wafer consistsof the region of the wafer between the surface of the silicon and adistance, D₁ (which is at least about 10 micrometers) as measured fromthe front surface and toward the central plane, the oxygen concentrationat a position within the denuded zone which is at a distance from thesilicon surface equal to one-half of D₁ will typically be at least about75% of the peak concentration of the interstitial oxygen concentrationanywhere in the denuded zone. For some oxygen precipitationheat-treatments, the interstitial oxygen concentration at this positionwill be even greater, i.e., at least 85%, 90% or even 95% of the maximumoxygen concentration anywhere in the denuded zone.

In a second embodiment of the present invention, a non-nitridingatmosphere is used instead of the nitriding atmosphere used in theheating (rapid thermal annealing) and cooling steps of the firstembodiment. Suitable non-nitriding atmospheres include argon, helium,neon, carbon dioxide, and other such non-oxidizing, non-nitridingelemental and compound gases, or mixtures of such gases. Thenon-nitriding atmosphere, like the nitriding atmosphere, may contain arelatively small partial pressure of oxygen, i.e., a partial pressureless than 0.01 atm. (10,000 ppma), more preferably less than 0.005 atm.(5,000 ppma), more preferably less than 0.002 atm. (2,000 ppma), andmost preferably less than 0.001 atm. (1,000 ppma).

In a third embodiment of the present invention, step S₁ (the thermaloxidation step) is omitted and the starting wafer has no more than anative oxide layer. When such a wafer is annealed in a nitrogenatmosphere, however, the effect differs from that which is observed whena wafer having an oxide layer which is greater in thickness than anative oxide layer (“enhanced oxide layer”) is annealed in nitrogen.When the wafer containing an enhanced oxide layer is annealed in anitrogen atmosphere, a substantially uniform increase in the vacancyconcentration is achieved throughout the wafer nearly, if notimmediately, upon reaching the annealing temperature; furthermore, thevacancy concentration does not appear to significantly increase as afunction of annealing time at a given annealing temperature. If thewafer lacks anything more than a native oxide layer and if the front andback surfaces of the wafer are annealed in nitrogen, however, theresulting wafer will have a vacancy concentration (number density)profile which is generally “U-shaped” for a cross-section of the wafer;that is, a maximum concentration will occur at or within severalmicrometers of the front and back surfaces and a relatively constant andlesser concentration will occur throughout the wafer bulk with theminimum concentration in the wafer bulk initially being approximatelyequal to the concentration which is obtained in wafers having anenhanced oxide layer. Furthermore, an increase in annealing time willresult in an increase in vacancy concentration in wafers lackinganything more than a native oxide layer.

Experimental evidence further suggests that this difference in behaviorfor wafers having no more than a native oxide layer and wafers having anenhanced oxide layer can be avoided by including molecular oxygen oranother oxidizing gas in the atmosphere. Stated another way, when wafershaving no more than a native oxide are annealed in a nitrogen atmospherecontaining a small partial pressure of oxygen, the wafer behaves thesame as wafers having an enhanced oxide layer. Without being bound toany theory, it appears that superficial oxide layers which are greaterin thickness than a native oxide layer serve as a shield which inhibitsnitridization of the silicon. This oxide layer may thus be present onthe starting wafer or formed, in situ, by growing an enhanced oxidelayer during the annealing step.

In accordance with the present invention, therefore, the atmosphereduring the rapid thermal annealing step preferably contains a partialpressure of at least about 0.0001 atm. (100 ppma), more preferably apartial pressure of at least about 0.0002 atm. (200 ppma). For thereasons previously discussed, however, the partial pressure of oxygenpreferably does not exceed 0.01 atm. (10,000 ppma), and is morepreferably less than 0.005 atm. (5,000 ppma), still more preferably lessthan 0.002 atm. (2,000 ppma), and most preferably less than 0.001 atm.(1,000 ppma).

In other embodiments of the present invention, the front and backsurfaces of the wafer may be exposed to different atmospheres, each ofwhich may contain one or more nitriding or non-nitriding gases. Forexample, the back surface of the wafer may be exposed to a nitridingatmosphere as the front surface is exposed to a non-nitridingatmosphere. Alternatively, multiple wafers (e.g., 2, 3 or more wafers)may be simultaneously annealed while being stacked in face-to-facearrangement; when annealed in this manner, the faces which are inface-to-face contact are mechanically shielded from the atmosphereduring the annealing. Alternatively, and depending upon the atmosphereemployed during the rapid thermal annealing step and the desired oxygenprecipitation profile of the wafer, the oxide layer may be formed onlyupon the side of the wafer at which the denuded zone is desired, e.g.,front surface 3 of the wafer (see FIG. 1).

The starting material for the process of the present invention may be apolished silicon wafer, or alternatively, a silicon wafer which has beenlapped and etched but not polished. In addition, the wafer may havevacancy or self-interstitial point defects as the predominant intrinsicpoint defect. For example, the wafer may be vacancy dominated fromcenter to edge, self-interstitial dominated from center to edge, or itmay contain a central core of vacancy of dominated material surroundedby an axially symmetric ring of self-interstitial dominated material.

More particularly, referring now to FIG. 15, such wafers may be obtainedfrom a single crystal silicon ingot 100, grown in accordance with theCzochralski method, comprising a central axis 120, a seed-cone 140, anend-cone 160 and a constant diameter portion 180 between the seed-coneand the end-cone. The constant diameter portion has a circumferentialedge 200 and a radius 40 extending from the central axis 120 to thecircumferential edge 200. In addition, the ingot may contain an axiallysymmetric region of self-interstitial dominated material 60, which issubstantially free of agglomerated intrinsic point defects, surroundinga generally cylindrical region of vacancy dominated material 80, aportion or all of which may also be substantially free of agglomeratedintrinsic point defects; alternatively, the self-interstitial dominatedregion 60 may extend from center to edge, or the vacancy dominatedregion 80 may extend from center to edge.

In one embodiment, an axially symmetric region 90 has a width, asmeasured along the radius 40 extending from V/I boundary 20 to axis 120,which is at least about 15 mm in width, and preferably has a width whichis at least about 7.5%, more preferably at least about 15%, still morepreferably at least about 25%, and most preferably at least about 50% ofthe radius of the constant diameter portion of the ingot. In aparticularly preferred embodiment, axially symmetric region 90 includesaxis 120 of the ingot, i.e., the axially symmetric region 90 and thegenerally cylindrical region 80 coincide. Stated another way, ingot 100includes a generally cylindrical region of vacancy dominated material80, at least a portion of which is free of agglomerated defects. Inaddition, axially symmetric region 90 extends over a length of at leastabout 20%, preferably at least about 40%, more preferably at least about60%, and still more preferably at least about 80% of the length of theconstant diameter portion of the ingot.

In another embodiment, such wafers may be obtained from a single crystalsilicon ingot 100 comprising an axially symmetric region 60 whichgenerally has a width 220, as measured from circumferential edge 200radially inward toward central axis 120, of at least about 30%, and insome embodiments at least about 40%, at least about 60%, or even atleast about 80% of the radius of the constant diameter portion of theingot. In addition, the axially symmetric region generally extends overa length 260 of at least about 20%, preferably at least about 40%, morepreferably at least about 60%, and still more preferably at least about80% of the length of the constant diameter portion of the ingot.

If an epitaxial layer is to be deposited upon an ideal precipitatingwafer, the process of the present invention may be carried out eitherbefore or after the epitaxial deposition. If carried out before, it maybe desirable to stabilize the oxygen precipitate nucleation centers inthe wafer after the process of the present invention and before theepitaxial deposition. If carried out after, it may be desirable to carryout the process of the present invention in the epitaxial reactorimmediately after the epitaxial deposition, provided the cooling ratesrequired by the process of the present invention can be achieved.

The measurement of crystal lattice vacancies in single crystal siliconcan be carried out by platinum diffusion analysis. In general, platinumis deposited on the samples and diffused in a horizontal surface withthe diffusion time and temperature preferably being selected such thatthe Frank-Turnbull mechanism dominates the platinum diffusion, but whichis sufficient to reach the steady-state of vacancy decoration byplatinum atoms. For wafers having vacancy concentrations which aretypical for the present invention, a diffusion time and temperature of730° C. for 20 minutes may be used, although more accurate trackingappears to be attainable at a lesser temperature, e.g., about 680° C. Inaddition, to minimize a possible influence by silicidation processes,the platinum deposition method preferably results in a surfaceconcentration of less than one monolayer. Platinum diffusion techniquesare described elsewhere, for example, by Jacob et al., J. Appl. Phys.,vol. 82, p. 182 (1997); Zimmermann and Ryssel, “The Modeling of PlatinumDiffusion In Silicon Under Non-Equilibrium Conditions,” J.Electrochemical Society, vol. 139, p. 256 (1992); Zimmermann, Goesele,Seilenthal and Eichiner, “Vacancy Concentration Wafer Mapping InSilicon,” Journal of Crystal Growth, vol. 129, p. 582 (1993); Zimmermannand Falster, “Investigation Of The Nucleation of Oxygen Precipitates inCzochralski Silicon At An Early Stage,” Appl. Phys. Lett., vol. 60, p.3250 (1992); and Zimmermann and Ryssel, Appl. Phys. A, vol. 55, p. 121(1992).

Preparation of Single Crystal Silicon Containing an Axially SymmetricRegion

Based upon experimental evidence to date, it appears that the type andinitial concentration of intrinsic point defects is initially determinedas the ingot cools from the temperature of solidification (i.e., about1410° C.) to a temperature greater than 1300° C. (i.e., at least about1325° C., at least about 1350° C. or even at least about 1375° C.). Thatis, the type and initial concentration of these defects are controlledby the ratio v/G₀, where v is the growth velocity and G₀ is the averageaxial temperature gradient over this temperature range.

Referring to FIG. 11, for increasing values of v/G₀, a transition fromdecreasingly self-interstitial dominated growth to increasingly vacancydominated growth occurs near a critical value of v/G₀ which, based uponcurrently available information, appears to be about 2.1×10⁻⁵ cm²/sK,where G₀ is determined under conditions in which the axial temperaturegradient is constant within the temperature range defined above. At thiscritical value, the concentrations of these intrinsic point defects areat equilibrium.

As the value of v/G₀ exceeds the critical value, the concentration ofvacancies increases. Likewise, as the value of v/G₀ falls below thecritical value, the concentration of self-interstitials increases. Ifthese concentrations reach a level of critical supersaturation in thesystem, and if the mobility of the point defects is sufficiently high, areaction, or an agglomeration event, will likely occur. Agglomeratedintrinsic point defects in silicon can severely impact the yieldpotential of the material in the production of complex and highlyintegrated circuits.

In accordance with the present invention, it has been discovered thatthe reactions in which vacancies within the silicon matrix react toproduce agglomerated vacancy defects and in which self-interstitialswithin the silicon matrix react to produce agglomerated interstitialdefects can be suppressed. Without being bound to any particular theory,it is believed that the concentration of vacancies andself-interstitials is controlled during the growth and cooling of thecrystal ingot in the process of the present invention, such that thechange in free energy of the system never exceeds a critical value atwhich the agglomeration reactions spontaneously occur to produceagglomerated vacancy or interstitial defects.

In general, the change in system free energy available to drive thereaction in which agglomerated vacancy defects are formed from vacancypoint defects or agglomerated interstitial defects are formed fromself-interstitial atoms in single crystal silicon is governed byEquation (1): $\begin{matrix}{{\Delta\quad G_{V/I}} = {k\quad T\quad{\ln\left( \frac{\left\lbrack {V/I} \right\rbrack}{\left\lfloor {V/I} \right\rfloor^{eq}} \right)}}} & (1)\end{matrix}$wherein

ΔG_(V/I) is the change in free energy for the reaction which formsagglomerated vacancy defects or the reaction which forms theinterstitial defects, as applicable,

k is the Boltzmann constant,

T is the temperature in K,

[V/I] is the concentration of vacancies or interstitials, as applicable,at a point in space and time in the single crystal silicon, and

[V/I]^(eq) is the equilibrium concentration of vacancies orinterstitials, as applicable, at the same point in space and time atwhich [V/I] occurs and at the temperature, T.

According to this equation, for a given concentration of vacancies, [V],a decrease in the temperature, T, generally results in an increase inΔG_(V)due to a sharp decrease in [V]^(eq) with temperature. Similarly,for a given concentration of interstitials, [I], a decrease in thetemperature, T, generally results in an increase in ΔG_(I) due to asharp decrease in [I]^(eq) with temperature.

FIG. 12 schematically illustrates the change in ΔG_(I) and theconcentration of silicon self-interstitials for an ingot which is cooledfrom the temperature of solidification without simultaneously employingsome means for suppression of the concentration of siliconself-interstitials. As the ingot cools, ΔG_(I) increases according toEquation (1), due to the increasing supersaturation of [I], and theenergy barrier for the formation of agglomerated interstitial defects isapproached. As cooling continues, this energy barrier is eventuallyexceeded, at which point a reaction occurs. This reaction results in theformation of agglomerated interstitial defects and the concomitantdecrease in ΔG_(I) as the supersaturated system is relaxed, i.e., as theconcentration of [I] decreases.

Similarly, as an ingot is cooled from the temperature of solidificationwithout simultaneously employing some means for suppression of theconcentration of vacancies, ΔG_(V) increases according to Equation (1),due to the increasing supersaturation of [V], and the energy barrier forthe formation of agglomerated vacancy defects is approached. As coolingcontinues, this energy barrier is eventually exceeded, at which point areaction occurs. This reaction results in the formation of agglomeratedvacancy defects and the concomitant decrease in ΔG_(V) as thesupersaturated system is relaxed.

The agglomeration of vacancies and interstitials can be avoided inregions of vacancy and interstitial dominated material, respectively, asthe ingot cools from the temperature of solidification by maintainingthe free energy of the vacancy system and the interstitial system atvalues which are less than those at which agglomeration reactions occur.In other words, the system can be controlled so as to never becomecritically supersaturated in vacancies or interstitials. This can beachieved by establishing initial concentrations of vacancies andinterstitials (controlled by v/G₀(r) as hereinafter defined) which aresufficiently low such that critical supersaturation is never achieved.In practice, however, such concentrations are difficult to achieveacross an entire crystal radius and, in general, therefore, criticalsupersaturation may be avoided by suppressing the initial vacancyconcentration and the initial interstitial concentration subsequent tocrystal solidification, i.e.,

subsequent to establishing the initial concentration determined byv/G₀(r).

Surprisingly, it has been found that due to the relatively largemobility of self-interstitials, which is generally about 10⁻⁴cm²/second, and to a lesser extent, to the mobility of vacancies, it ispossible to effect the suppression of interstitials and vacancies overrelatively large distances, i.e. distances of about 5 cm to about 10 cmor more, by the radial diffusion of self-interstitials to sinks locatedat the crystal surface or to vacancy dominated regions located withinthe crystal. Radial diffusion can be effectively used to suppress theconcentration of self-interstitials and vacancies, provided sufficienttime is allowed for the radial diffusion of the initial concentration ofintrinsic point defects. In general, the diffusion time will depend uponthe radial variation in the initial concentration of self-interstitialsand vacancies, with lesser radial variations requiring shorter diffusiontimes.

Typically, the average axial temperature gradient, G₀, increases as afunction of increasing radius for single crystal silicon, which is grownaccording to the Czochralski method. This means that the value of v/G₀is typically not singular across the radius of an ingot. As a result ofthis variation, the type and initial concentration of intrinsic pointdefects is not constant. If the critical value of v/G₀, denoted in FIGS.13 and 14 as the V/I boundary 20, is reached at some point along theradius 40 of the ingot, the material will switch from being vacancydominated to self-interstitial dominated. In addition, the ingot willcontain an axially symmetric region of self-interstitial dominatedmaterial 60 (in which the initial concentration of siliconself-interstitial atoms increases as a function of increasing radius),surrounding a generally cylindrical region of vacancy dominated material80 (in which the initial concentration of vacancies decreases as afunction of increasing radius).

As an ingot containing a V/I boundary is cooled from the temperature ofsolidification, radial diffusion of interstitial atoms and vacanciescauses a radially inward shift in the V/I boundary due to arecombination of self-interstitials with vacancies. In addition, radialdiffusion of self-interstitials to the surface of the crystal will occuras the crystal cools. The surface of the crystal is capable ofmaintaining near equilibrium point defect concentrations as the crystalcools. Radial diffusion of point defects will tend to reduce theself-interstitial concentration outside the V/I boundary and the vacancyconcentration inside the V/I boundary. If enough time is allowed fordiffusion, therefore, the concentration of vacancy and interstitialseverywhere may be such that ΔG_(V) and ΔG_(I) will be less than thecritical values at which the vacancy agglomeration reaction and theinterstitial agglomeration reactions occur.

Referring again to FIG. 15, the crystal growth conditions, includinggrowth velocity, v, the average axial temperature gradient, G₀, and thecooling rate are preferably controlled to cause the formation of asingle crystal silicon ingot 100, as previously described. Preferably,these conditions are controlled to cause the formation of an axiallysymmetric region of interstitial dominated material 60 and a generallycylindrical region of vacancy dominated material 80, which may or maynot contain an axially symmetric region of agglomerated intrinsic pointdefect-free material 90. When present, the axially symmetric region 90has a width which may vary, as previously discussed; likewise, whenpresent, the axially symmetric region 60 also has a width which mayvary, as previously discussed.

The width of axially symmetric regions 60 and 90 may have some variationalong the length of the central axis 120. For an axially symmetricregion of a given length, therefore, the width 220 of axially symmetricregion 60 is determined by measuring the distance from thecircumferential edge 200 of the ingot 100 radially toward a point whichis farthest from the central axis. In other words, the width is measuredsuch that the minimum distance within the given length of the axiallysymmetric region 60 is determined. Similarly, the width of axiallysymmetric region 90 is determined by measuring the distance from the V/Iboundary 20 radially toward a point which is farthest from the centralaxis 120. In other words, the width is measured such that the minimumdistance within the given length of the axially symmetric region 90 isdetermined.

The growth velocity, v, and the average axial temperature gradient, G₀,(as previously defined) are typically controlled such that the ratiov/G₀ ranges in value from about 0.5 to about 2.5 times the criticalvalue of v/G₀ (i.e., about 1×10⁻⁵ cm²/sK to about 5×10⁻⁵ cm²/sK basedupon currently available information for the critical value of v/G₀).Preferably, the ratio v/G₀ will range in value from about 0.6 to about1.5 times the critical value of v/G₀ (i.e., about 1.3×10⁻⁵ cm²/sK toabout 3×10⁻⁵ cm²/sK based upon currently available information for thecritical value of v/G₀). Most preferably, the ratio v/G₀ will range invalue from about 0.75 to about 1.25 times the critical value of v/G₀(i.e., about 1.6×10⁻⁵ cm²/sK to about 2.1×10⁻⁵ cm²/sK based uponcurrently available information for the critical value of v/G₀). In oneparticularly preferred embodiment, v/G₀ within the axially symmetricregion 80 has a value falling between the critical value of v/G₀ and 1.1times the critical value of v/G₀. In another particularly preferredembodiment, v/G₀ within the axially symmetric region 60 has a valuefalling between about 0.75 times the critical value of v/G₀ and thecritical value of v/G₀.

To maximize the width of the axially symmetric regions 60 or 90, it ispreferred that the ingot be cooled from the solidification temperatureto a temperature in excess of about 1050° C. over a period of (i) atleast about 5 hours, preferably at least about 10 hours, and morepreferably at least about 15 hours for 150 mm nominal diameter siliconcrystals, (ii) at least about 5 hours, preferably at least about 10hours, more preferably at least about 20 hours, still more preferably atleast about 25 hours, and most preferably at least about 30 hours for200 mm nominal diameter silicon crystals, and (iii) at least about 20hours, preferably at least about 40 hours, more preferably at leastabout 60 hours, and most preferably at least about 75 hours for siliconcrystals having a nominal diameter greater than 200 mm. Control of thecooling rate can be achieved by using any means currently known in theart for minimizing heat transfer, including the use of insulators,heaters, radiation shields, and magnetic fields.

Control of the average axial temperature gradient, G₀, may be achievedthrough the design of the “hot zone” of the crystal puller, i.e. thegraphite (or other materials) that makes up the heater, insulation, heatand radiation shields, among other things. Although the designparticulars may vary depending upon the make and model of the crystalpuller, in general, G₀ may be controlled using any of the meanscurrently known in the art for controlling heat transfer at themelt/solid interface, including reflectors, radiation shields, purgetubes, light pipes, and heaters. In general, radial variations in G₀ areminimized by positioning such an apparatus within about one crystaldiameter above the melt/solid interface. G₀ can be controlled further byadjusting the position of the apparatus relative to the melt andcrystal. This is accomplished either by adjusting the position of theapparatus in the hot zone, or by adjusting the position of the meltsurface in the hot zone. In addition, when a heater is employed, G₀ maybe further controlled by adjusting the power supplied to the heater.Any, or all, of these methods can be used during a batch Czochralskiprocess in which melt volume is depleted during the process.

It is generally preferred for some embodiments of the present inventionthat the average axial temperature gradient, G₀, be relatively constantas a function of diameter of the ingot. However, it should be noted thatas improvements in hot zone design allow for variations in G₀ to beminimized, mechanical issues associated with maintaining a constantgrowth rate become an increasingly important factor. This is because thegrowth process becomes much more sensitive to any variation in the pullrate, which in turn directly effects the growth rate, v. In terms ofprocess control, this means that it is favorable to have values for G₀which differ over the radius of the ingot. Significant differences inthe value of G₀, however, can result in a large concentration ofself-interstitials generally increasing toward the wafer edge and,thereby, increase the difficultly in avoiding the formation ofagglomerated intrinsic point defects.

In view of the foregoing, the control of G₀ involves a balance betweenminimizing radial variations in G₀ and maintaining favorable processcontrol conditions. Typically, therefore, the pull rate after about onediameter of the crystal length will range from about 0.2 mm/minute toabout 0.8 mm/minute. Preferably, the pull rate will range from about0.25 mm/minute to about 0.6 mm/minute and, more preferably, from about0.3 mm/minute to about 0.5 mm/minute. It is to be noted that the pullrate is dependent upon both the crystal diameter and crystal pullerdesign. The stated ranges are typical for 200 mm diameter crystals. Ingeneral, the pull rate will decrease as the crystal diameter increases.However, the crystal puller may be designed to allow pull rates inexcess of those stated here. As a result, most preferably the crystalpuller will be designed to enable the pull rate to be as fast aspossible while still allowing for the formation of an axially symmetricregion or regions in accordance with the present invention.

The amount of self-interstitial diffusion is controlled by controllingthe cooling rate as the ingot is cooled from the solidificationtemperature (about 1410° C.) to the temperature at which siliconself-interstitials become immobile, for commercially practical purposes.Silicon self-interstitials appear to be extremely mobile at temperaturesnear the solidification temperature of silicon, i.e. about 1410° C. Thismobility, however, decreases as the temperature of the single crystalsilicon ingot decreases. Generally, the diffusion rate ofself-interstitials slows such a considerable degree that they areessentially immobile for commercially practical time periods attemperatures less than about 700° C., and perhaps at temperatures asgreat as 800° C., 900° C., 1000° C., or even 1050° C.

It is to be noted in this regard that, although the temperature at whicha self-interstitial agglomeration reaction occurs may in theory varyover a wide range of temperatures, as a practical matter this rangeappears to be relatively narrow for conventional, Czochralski grownsilicon. This is a consequence of the relatively narrow range of initialself-interstitial concentrations which are typically obtained in silicongrown according to the Czochralski method. In general, therefore, aself-interstitial agglomeration reaction may occur, if at all, attemperatures within the range of about 1100° C. to about 800° C., andtypically at a temperature of about 1050° C.

Within the range of temperatures at which self-interstitials appear tobe mobile, and depending upon the temperature in the hot zone, thecooling rate will typically range from about 0.1° C./minute to about 3°C./minute. Preferably, the cooling rate will range from about 0.1°C./minute to about 1.5° C./minute, more preferably from about 0.1°C./minute to about 1° C./minute, and still more preferably from about0.1° C./minute to about 0.5° C./minute.

By controlling the cooling rate of the ingot within a range oftemperatures in which self-interstitials appear to be mobile, theself-interstitials may be given more time to diffuse to sinks located atthe crystal surface, or to vacancy dominated regions, where they may beannihilated. The concentration of such interstitials may therefore besuppressed, which act to prevent an agglomeration event from occurring.Utilizing the diffusivity of interstitials by controlling the coolingrate acts to relax the otherwise stringent v/G₀ requirements that may berequired in order to obtain an axially symmetric region free ofagglomerated defects. Stated another way, as a result of the fact thatthe cooling rate may be controlled in order to allow interstitials moretime to diffuse, a large range of v/G₀ values, relative to the criticalvalue, are acceptable for purposes of obtaining an axially symmetricregion free of agglomerated defects.

To achieve such cooling rates over appreciable lengths of the constantdiameter portion of the crystal, consideration must also be given to thegrowth process of the end-cone of the ingot, as well as the treatment ofthe ingot once end-cone growth is complete. Typically, upon completionof the growth of the constant diameter portion of the ingot, the pullrate will be increased in order to begin the tapering necessary to formthe end-cone. However, such an increase in pull rate will result in thelower segment of the constant diameter portion cooling more quicklywithin the temperature range in which interstitials are sufficientlymobile, as discussed above, As a result, these interstitials may nothave sufficient time to diffuse to sinks to be annihilated; that is, theconcentration in this lower segment may not be suppressed to asufficient degree and agglomeration of interstitial defects may result.

In order to prevent the formation of such defects from occurring in thislower segment of the ingot, it is therefore preferred that constantdiameter portion of the ingot have a uniform thermal history inaccordance with the Czochralski method. A uniform thermal history may beachieved by pulling the ingot from the silicon melt at a relativelyconstant rate during the growth of not only the constant diameterportion, but also during the growth of the end-cone of the crystal andpossibly subsequent to growth of the end-cone. The relatively constantrate may be achieved, for example, by (i) reducing the rates of rotationof the crucible and crystal during the growth of the end-cone relativeto the crucible and crystal rotation rates during the growth of theconstant diameter portion of the crystal, and/or (ii) increasing thepower supplied to the heater used to heat the silicon melt during thegrowth of the end-cone relative to the power conventionally suppliedduring end-cone growth. These additional adjustments of the processvariables may occur either individually or in combination.

When the growth of the end-cone is initiated, a pull rate for theend-cone is established such that, any segment of the constant diameterportion of the ingot which remains at a temperature in excess of about1050° C. experiences the same thermal history as other segment(s) of theconstant diameter portion of the ingot which contain an axiallysymmetric region free of agglomerated intrinsic point defects which havealready cooled to a temperature of less than about 1050° C.

As previously noted, a minimum radius of the vacancy dominated regionexists for which the suppression of agglomerated interstitial defectsmay be achieved. The value of the minimum radius depends on v/G₀(r) andthe cooling rate. As crystal puller and hot zone designs will vary, theranges presented above for v/G₀(r), pull rate, and cooling rate willalso vary. Likewise these conditions may vary along the length of agrowing crystal. Also as noted above, the width of the interstitialdominated region free of agglomerated interstitial defects is preferablymaximized. Thus, it is desirable to maintain the width of this region toa value which is as close as possible to, without exceeding, thedifference between the crystal radius and the minimum radius of thevacancy dominated region along the length of the growing crystal in agiven crystal puller.

The optimum width of axially symmetric regions 60 and 90 and therequired optimal crystal pulling rate profile for a given crystal pullerhot zone design may be determined empirically. Generally speaking, thisempirical approach involves first obtaining readily available data onthe axial temperature profile for an ingot grown in a particular crystalpuller, as well as the radial variations in the average axialtemperature gradient for an ingot grown in the same puller.Collectively, this data is used to pull one or more single crystalsilicon ingots, which are then analyzed for the presence of agglomeratedinterstitial defects. In this way, an optimum pull rate profile can bedetermined.

FIG. 16 is an image produced by a scan of the minority carrier lifetimeof an axial cut of a section of a 200 mm diameter ingot following aseries of oxygen precipitation heat-treatments which reveal defectdistribution patterns. It depicts an example in which a near-optimumpull rate profile is employed for a given crystal puller hot zonedesign. In this example, a transition occurs from a v/G₀(r) at which themaximum width of the interstitial dominated region is exceeded(resulting in the generation of regions of agglomerated interstitialdefects 280) to an optimum v/G₀(r) at which the axially symmetric regionhas the maximum width.

In addition to the radial variations in v/G₀ resulting from an increasein G₀ over the radius of the ingot, v/G₀ may also vary axially as aresult of a change in v, or as a result of natural variations in G₀ dueto the Czochralski process. For a standard Czochralski process, v isaltered as the pull rate is adjusted throughout the growth cycle, inorder to maintain the ingot at a constant diameter. These adjustments,or changes, in the pull rate in turn cause v/G₀ to vary over the lengthof the constant diameter portion of the ingot. In accordance with theprocess of the present invention, the pull rate is therefore controlledin order to maximize the width of the axially symmetric region of theingot. As a result, however, variations in the radius of the ingot mayoccur. In order to ensure that the resulting ingot has a constantdiameter, the ingot is therefore preferably grown to a diameter largerthan that which is desired. The ingot is then subjected to processesstandard in the art to remove excess material from the surface, thusensuring that an ingot having a constant diameter portion is obtained.

In general, it is easier to make vacancy dominated material free ofagglomerated defects when radial variation of the axial temperaturegradient, G₀(r), is minimized. Referring to FIG. 35, axial temperatureprofile for four separate hot zone configurations are illustrated. FIG.34 presents the variation in the axial temperature gradient, G₀(r), fromthe center of the crystal to one-half of the crystal radius, determinedby averaging the gradient from the solidification temperature to thetemperature indicated on the x-axis. When crystals were pulled in hotzones designated as Ver. 1 and Ver. 4 which have larger radial variationin G₀(r), it was not possible to obtain crystals having vacancydominated material from center to edge of any axial length which wasfree of agglomerated defects. When crystals were pulled in hot zonesdesignated as Ver. 2 and Ver. 3 which have lesser radial variation inG₀(r), however, it was possible to obtain crystals having vacancydominated material from center to edge which was free of agglomerateddefects for some axial length of the crystal.

In one embodiment of the process of the present invention, the initialconcentration of silicon self-interstitial atoms is controlled in theaxially symmetric, self-interstitial dominated region 60 of ingot 100.Referring again to FIG. 10, in general, the initial concentration ofsilicon self-interstitial atoms is controlled by controlling the crystalgrowth velocity, v, and the average axial temperature gradient, G₀, suchthat the value of the ratio v/G₀ is relatively near the critical valueof this ratio, at which the V/I boundary occurs. In addition, theaverage axial temperature gradient, G₀, can be established such that thevariation of G₀, i.e. G₀(r), (and thus, v/G₀(r)) as a the variation ofG₀ (and thus, v/G₀) as a function of the ingot radius is alsocontrolled.

In another embodiment of the present invention, V/G₀ is controlled suchthat no V/I boundary exists along the radius for at least a portion ofthe length of the ingot. In this length, the silicon is vacancydominated from center to circumferential edge and agglomerated vacancydefects are avoided in an axially symmetric region extending radiallyinward from the circumferential edge of the ingot principally bycontrolling V/G₀. That is, the growth conditions are controlled so thatv/G₀ has a value falling between the critical value of v/G₀ and 1.1times the critical value of v/G₀.

Visual Detection of Agglomerated Defects

Agglomerated defects may be detected by a number of differenttechniques. For example, flow pattern defects, or D-defects, aretypically detected by preferentially etching the single crystal siliconsample in a Secco etch solution for about 30 minutes, and thensubjecting the sample to microscopic inspection. (see, e.g., H.Yamagishi et al., Semicond. Sci. Technol. 7, A135 (1992)). Althoughstandard for the detection of agglomerated vacancy defects, this processmay also be used to detect agglomerated interstitial defects. When thistechnique is used, such defects appear as large pits on the surface ofthe sample when present.

Agglomerated defects may also be detected using laser scatteringtechniques, such as laser scattering tomography, which typically have alower defect density detection limit that other etching techniques.

Additionally, agglomerated intrinsic point defects may be visuallydetected by decorating these defects with a metal capable of diffusinginto the single crystal silicon matrix upon the application of heat.Specifically, single crystal silicon samples, such as wafers, slugs orslabs, may be visually inspected for the presence of such defects byfirst coating a surface of the sample with a composition containing ametal capable of decorating these defects, such as a concentratedsolution of copper nitrate. The coated sample is then heated to atemperature between about 900° C. and about 1000° C. for about 5 minutesto about 15 minutes in order to diffuse the metal into the sample. Theheat treated sample is then cooled to room temperature, thus causing themetal to become critically supersaturated and precipitate at siteswithin the sample matrix at which defects are present.

After cooling, the sample is first subjected to a non-defect delineatingetch, in order to remove surface residue and precipitants, by treatingthe sample with a bright etch solution for about 8 to about 12 minutes.A typical bright etch solution comprises about 55 percent nitric acid(70% solution by weight), about 20 percent hydrofluoric acid (49%solution by weight), and about 25 percent hydrochloric acid(concentrated solution).

The sample is then rinsed with deionized water and subjected to a secondetching step by immersing the sample in, or treating it with, a Secco orWright etch solution for about 35 to about 55 minutes. Typically, thesample will be etched using a Secco etch solution comprising about a 1:2ratio of 0.15 M potassium dichromate and hydrofluoric acid (49% solutionby weight). This etching step acts to reveal, or delineate, agglomerateddefects which may be present.

In general, regions of interstitial and vacancy dominated material freeof agglomerated defects can be distinguished from each other and frommaterial containing agglomerated defects by the copper decorationtechnique described above. Regions of defect-free interstitial dominatedmaterial contain no decorated features revealed by the etching whereasregions of defect-free vacancy dominated material (prior to ahigh-temperature oxygen nuclei dissolution treatment as described above)contain small etch pits due to copper decoration of the oxygen nuclei.

Definitions

As used herein, the following phrases or terms shall have the givenmeanings: “agglomerated intrinsic point defects” mean defects caused (i)by the reaction in which vacancies agglomerate to produce D-defects,flow pattern defects, gate oxide integrity defects, crystal originatedparticle defects, crystal originated light point defects, and other suchvacancy related defects, or (ii) by the reaction in whichself-interstitials agglomerate to produce dislocation loops andnetworks, and other such self-interstitial related defects;“agglomerated interstitial defects” shall mean agglomerated intrinsicpoint defects caused by the reaction in which silicon self-interstitialatoms agglomerate; “agglomerated vacancy defects” shall meanagglomerated vacancy point defects caused by the reaction in whichcrystal lattice vacancies agglomerate; “radius” means the distancemeasured from a central axis to a circumferential edge of a wafer oringot; “substantially free of agglomerated intrinsic point defects”shall mean a concentration of agglomerated defects which is less thanthe detection limit of these defects, which is currently about 10³defects/cm³; “V/I boundary” means the position along the radius of aningot or wafer at which the material changes from vacancy dominated toself-interstitial dominated; and “vacancy dominated” and“self-interstitial dominated” mean material in which the intrinsic pointdefects are predominantly vacancies or self-interstitials, respectively.

EXAMPLES

Examples 1 through 4 illustrate the ideal oxygen precipitation processof the present invention. Examples 5 through 11 illustrate thepreparation of single crystal silicon containing an axially symmetricregion of vacancy dominated material, self-interstitial dominatedmaterial, or both, which is substantially free of agglomerated intrinsicpoint defects, as previously discussed. All of these Examples shouldtherefore not be interpreted in a limiting sense.

Ideal Oxygen Precipitation Process Example 1

Silicon single crystals were pulled by the Czochralski method, slicedand polished to form silicon wafers. These wafers were then subjected toa surface oxidation step (S₁), rapid thermal annealing step in nitrogenor argon (S₂), rapidly cooled (S₃), and subjected to an oxygenstabilization and growth step (S₄) under the conditions set forth inTable I. The initial oxygen concentration of the wafers (O_(i)) beforesteps S₁-S₄, the oxygen precipitate density in the bulk of the wafersafter step S₄ (OPD), and the depth of the denuded zone after step S₄(DZ) are also reported in Table I. TABLE I Sample 4-7 4-8 3-14 S₁ 15 minat 15 min at none 1,000° C. 1,000° C. in N₂ + ˜1% in N₂ + ˜1% O₂ O₂ S₂35 seconds 35 seconds 35 seconds at 1250° C. at 1250° C. at 1250° C. inN₂ in Ar in N₂ S₃ 100° C./sec 100° C./sec 100° C./sec S₄ 4 hr at 4 hr at4 hr at 800° C. + 16 hr 800° C. + 16 hr 800° C. + 16 hr at at at 1,000°C. 1,000° C. 1,000° C. in N₂ in N₂ in N₂ O_(i) 7 × 10¹⁷ 6.67 × 10¹⁷  7.2× 10¹⁷ (atoms/cm³) OPD 1 × 10¹⁰  4.4 × 10⁹ 1.69 × 10¹⁰ (atoms/cm³) DZ 7095 0 (depth in μm)

FIGS. 2, 3, and 4 show cross-sections of the resulting wafers (thesefigures are enlargements of photographs taken at a magnification of200×); sample 4-7 is shown in FIG. 2, sample 4-8 is shown in FIG. 3, andsample 3-14 is shown in FIG. 4.

In addition, the concentration of crystal lattice vacancies in thesample 4-7 was mapped using a platinum diffusion technique. A plot ofplatinum concentration versus depth from the surface of the wafer (adepth of 0 micrometers corresponding to the front side of the wafer)appears in FIG. 5.

Example 2

To demonstrate that the process of the present invention is relativelyindependent of oxygen concentration for Czochralski-grown siliconwafers, three wafers having different oxygen concentrations weresubjected to the same series of steps described in Example 1. Theconditions for each of these steps, the initial oxygen concentration ofthe wafers (O_(i)) before steps S₁-S₄, the oxygen precipitate density(OPD) in the bulk of the wafers after step S₄, and the depth of thedenuded zone (DZ) after step S₄ as measured from the surface of thewafer are reported in Table II. FIGS. 6, 7, and 8 show cross-sections ofthe resulting wafers (these figures are enlargements of photographstaken at a magnification of 200×); sample 3-4 is shown in FIG. 6, sampleis shown in FIG. 7, and sample 3-6 is shown in FIG. 8. TABLE II Sample3-4 3-5 3-6 S₁ 15 min at 15 min at 15 min at 1,000° C. 1,000° C. 1,000°C. in N₂ + ˜1% in N₂ + ˜1% in N₂ + ˜1% O₂ O₂ O₂ S₂ 35 seconds 35 seconds35 seconds at 1250° C. at 1250° C. at 1250° C. in N₂ in N₂ in N₂ S₃ 125°C./sec 125° C./sec 125° C./sec S₄ 4 hr at 4 hr at 4 hr at 800° C. + 16hr 800° C. + 16 hr 800° C. + 16 hr at at at 1,000° C. 1,000° C. 1,000°C. in N₂ in N₂ in N₂ O_(i) 6 × 10¹⁷ 7 × 10¹⁷ 8 × 10¹⁷ (atoms/cm³) OPD 4× 10¹⁰ 1 × 10¹⁰ 6 × 10¹⁰ (atoms/cm³) DZ ˜40 ˜40 ˜40 (depth in μm)

Example 3

To demonstrate that the process of the present invention was relativelyindependent of the conditions used for the oxygen precipitatestabilization and growth step (S₄), a wafer (sample 1-8) having the sameinitial oxygen concentration was subjected to the same series of stepsdescribed in Example 2 for sample 3-4 except that a proprietary,commercial 16 Mb DRAM process was used as the oxygen precipitatestabilization and growth step (S₄). FIG. 9 shows a cross-section of theresulting wafer (this figure is an enlargement of a photograph taken ata magnification of 200×). After step S₄, samples 1-8 and 3-4 hadcomparable bulk oxygen precipitate densities (7×10¹⁰/cm³ for sample 1-8versus 4×10¹⁰/cm³ for sample 3-4) and comparable denuded zone depths(approximately 40 micrometers).

Example 4

This example illustrates the trend that may be observed in the densityof bulk microdefects (BMD), i.e., the density of oxygen precipitants,and the depth of the denuded zone (DZ) resulting from an increase in theconcentration of oxygen in the atmosphere during the heat-treatment.Three different sets of wafers were subjected to rapid thermal annealingunder varying process conditions. The wafers in Set A were annealed at1200° C. for 30 seconds under a nitrogen atmosphere; the wafers in Set Bwere annealed under the same conditions for 20 seconds; and, the wafersin Set C were annealed at 1200° C. for 30 seconds under an argonatmosphere. A pre-oxidation step was not performed on any of the wafersin the three sets in this Example.

As indicated by Table III, below, the oxygen partial pressure wasincreased for each wafer within a given Set. Once annealing wascompleted, the BMD density and DZ depth for each wafer was determined bymeans standard in the art. The results are present in Table III, below.TABLE III Wafer Oxygen Partial BMD Density DZ Depth Set Pressure(defects/cm⁻³) (microns) A 250 ppma 6.14 × 10⁹ 70 A 500 ppma 6.24 × 10⁹80 A 1000 ppma 2.97 × 10⁹ 80 A 2000 ppma 7.02 × 10⁸ 100  A 5000 ppma2.99 × 10⁷ ND A 1 × 10⁶ ppma 6.03 × 10⁶ ND B 500 ppma 2.59 × 10⁹ 80 B1000 ppma 1.72 × 10⁹ 100  B 2000 ppma 9.15 × 10⁸ 100  B 5000 ppma 2.65 ×10⁷ ND B 1 × 10⁶ ppma 2.17 × 10⁶ ND C 250 ppma 2.65 × 10⁹ 90 C 500 ppma4.03 × 10⁹ 70 C 1000 ppma 1.72 × 10⁹ 140  C 5000 ppma 1.69 × 10⁸ 120 ND = not determined

The above data shows that as the partial pressure of oxygen in theatmosphere increases, the number density of bulk microdefects decreases.In addition, when the oxygen partial pressure reaches 10,000 ppma, thenumber density of bulk microdefects is indistinguishable from the numberdensity of bulk microdefects which is observed in wafers which have beensubjected to an oxygen precipitation heat-treatment without a priorrapid thermal annealing in accordance with the present invention.

Single Crystal Silicon Containing an Axially Symmetric Region Example 5Optimization Procedure For A Crystal Puller Having A Pre-existing HotZone Design

A first 200 mm single crystal silicon ingot was grown under conditionsin which the pull rate was ramped linearly from about 0.75 mm/min. toabout 0.35 mm/min. over the length of the crystal. FIG. 17 shows thepull rate as a function of crystal length. Taking into account thepre-established axial temperature profile of a growing 200 mm ingot inthe crystal puller and the pre-established radial variations in theaverage axial temperature gradient, G₀, i.e., the axial temperaturegradient at the melt/solid interface, these pull rates were selected toinsure that ingot would be vacancy dominated material from the center tothe edge at one end of the ingot and interstitial dominated materialfrom the center to the edge of the other end of the ingot. The growningot was sliced longitudinally and analyzed to determine where theformation of agglomerated interstitial defects begins.

FIG. 18 is an image produced by a scan of the minority carrier lifetimeof an axial cut of the ingot over a section ranging from about 635 mm toabout 760 mm from the shoulder of the ingot following a series of oxygenprecipitation heat-treatments which reveal defect distribution patterns.At a crystal position of about 680 mm, a band of agglomeratedinterstitial defects 280 can be seen. This position corresponds to acritical pull rate of v* (680 mm)=0.33 mm/min. At this point, the widthof the axially symmetric region 60 (a region which is interstitialdominated material but which lacks agglomerated interstitial defects) isat its maximum; the width of the vacancy dominated region 80, R_(v)*(680) is about 35 mm and the width of the axially symmetric region,R_(I)* (680) is about 65 mm.

A series of four single crystal silicon ingots were then grown at steadystate pull rates which were somewhat greater than and somewhat less thanthe pull rate at which the maximum width of the axially symmetric regionof the first 200 mm ingot was obtained. FIG. 19 shows the pull rate as afunction of crystal length for each of the four crystals, labeled,respectively, as 1-4. These four crystals were then analyzed todetermine the axial position (and corresponding pull rate) at whichagglomerated interstitial defects first appear or disappear. These fourempirically determined points (marked “*”) are shown in FIG. 19.Interpolation between and extrapolation from these points yielded acurve, labeled v* (Z) in FIG. 19. This curve represents, to a firstapproximation, the pull rate for 200 mm crystals as a function of lengthin the crystal puller at which the axially symmetric region is at itsmaximum width.

Growth of additional crystals at other pull rates and further analysisof these crystals would further refine the empirical definition of v*(Z).

Example 6 Reduction of Radial Variation in G₀(r)

FIGS. 20 and 21 illustrate the improvement in quality that can beachieved by reduction of the radial variation in the axial temperaturegradient at the melt/solid interface, G₀(r). The initial concentration(about 1 cm from the melt/solid interface) of vacancies andinterstitials are calculated for two cases with different G₀(r): (1)G₀(r)=2.65+5×10⁻⁴r² (K/mm) and (2) G₀(r)=2.65+5×10⁻⁵r² (K/mm). For eachcase the pull rate was adjusted such that the boundary betweenvacancy-rich silicon and interstitial-rich silicon is at a radius of 3cm. The pull rate used for case 1 and 2 were 0.4 and 0.35 mm/min,respectively. From FIG. 21 it is clear that the initial concentration ofinterstitials in the interstitial-rich portion of the crystal isdramatically reduced as the radial variation in the initial axialtemperature gradient is reduced. This leads to an improvement in thequality of the material since it becomes easier to avoid the formationof interstitial defect clusters due to supersaturation of interstitials.

Example 7 Increased Out-Diffusion Time for Interstitials

FIGS. 22 and 23 illustrate the improvement in quality that can beachieved by increasing the time for out-diffusion of interstitials. Theconcentration of interstitials is calculated for two cases withdiffering axial temperature profiles in the crystal, dT/dz. The axialtemperature gradient at the melt/solid interface is the same for bothcases, so that the initial concentration (about 1 cm from the melt/solidinterface) of interstitials is the same for both cases. In this example,the pull rate was adjusted such that the entire crystal isinterstitial-rich. The pull rate was the same for both cases, 0.32mm/min. The longer time for interstitial out-diffusion in case 2 resultsin an overall reduction of the interstitial concentration. This leads toan improvement in the quality of the material since it becomes easier toavoid the formation of interstitial defect clusters due tosupersaturation of interstitials.

Example 8

A 700 mm long, 150 mm diameter crystal was grown with a varying pullrate. The pull rate varied nearly linearly from about 1.2 mm/min at theshoulder to about 0.4 mm/min at 430 mm from the shoulder, and thennearly linearly back to about 0.65 mm/min at 700 mm from the shoulder.Under these conditions in this particular crystal puller, the entireradius is grown under interstitial-rich conditions over the length ofcrystal ranging from about 320 mm to about 525 mm from the shoulder ofthe crystal. Referring to FIG. 24, at an axial position of about 525 mmand a pull rate of about 0.47 mm/min, the crystal is free ofagglomerated intrinsic point defects clusters across the entirediameter. Stated another way, there is one small section of the crystalin which the width of the axially symmetric region, i.e., the regionwhich is substantially free of agglomerated defects, is equal to theradius of the ingot.

Example 9

As described in Example 5, a series of single crystal silicon ingotswere grown at varying pull rates and then analyzed to determine theaxial position (and corresponding pull rate) at which agglomeratedinterstitial defects first appeared or disappeared. Interpolationbetween and extrapolation from these points, plotted on a graph of pullrate v. axial position, yielded a curve which represents, to a firstapproximation, the pull rate for a 200 mm crystal as a function oflength in the crystal puller at which the axially symmetric region is atits maximum width. Additional crystals were then grown at other pullrates and further analysis of these crystals was used to refine thisempirically determined optimum pull rate profile.

Using this data and following this optimum pull rate profile, a crystalof about 1000 mm in length and about 200 mm in diameter was grown.Slices of the grown crystal, obtained from various axial position, werethen analyzed using oxygen precipitation methods standard in the art inorder to (i) determine if agglomerated interstitial defects were formed,

and (ii) determine, as a function of the radius of the slice, theposition of the V/I boundary. In this way the presence of an axiallysymmetric region was determined, as well as the width of this region afunction of crystal length or position.

The results obtained for axial positions ranging from about 200 mm toabout 950 mm from the shoulder of the ingot are present in the graph ofFIG. 25. These results show that a pull rate profile may be determinedfor the growth of a single crystal silicon ingot such that the constantdiameter portion of the ingot may contain an axially symmetric regionhaving a width, as measured from the circumferential edge radiallytoward the central axis of the ingot, which is at least about 40% thelength of the radius of the constant diameter portion. In addition,these results show that this axially symmetric region may have a length,as measured along the central axis of the ingot, which is about 75% ofthe length of the constant diameter portion of the ingot.

Example 10

A single crystal silicon ingot have a length of about 1100 mm and adiameter of about 150 mm was grown with a decreasing pull rate. The pullrate at the shoulder of the constant diameter portion of the ingot wasabout 1 mm/min. The pull rate decreased exponentially to about 0.4mm/min., which corresponded to an axial position of about 200 mm fromthe shoulder. The pull rate then decreased linearly until a rate ofabout 0.3 mm/min. was reached near the end of the constant diameterportion of the ingot.

Under these process conditions in this particular hot zoneconfiguration, the resulting ingot contains a region wherein the axiallysymmetric region has a width which about equal to the radius of theingot. Referring now to FIGS. 26 a and 26 b, which are images producedby a scan of the minority carrier lifetime of an axial cut of a portionof the ingot following a series of oxygen precipitation heat treatments,consecutive segments of the ingot, ranging in axial position from about100 mm to about 250 mm and about 250 mm to about 400 mm are present. Itcan be seen from these figures that a region exists within the ingot,ranging in axial position from about 170 mm to about 290 mm from theshoulder, which is free of agglomerated intrinsic point defects acrossthe entire diameter. Stated another way, a region is present within theingot wherein the width of the axially symmetric region, i.e., theregion which is substantially free of agglomerated interstitial defects,is about equal to the radius of the ingot.

In addition, in a region ranging from an axially position from about 125mm to about 170 mm and from about 290 mm to greater than 400 mm thereare axially symmetric regions of interstitial dominated material free ofagglomerated intrinsic point defects surrounding a generally cylindricalcore of vacancy dominated material which is also free of agglomeratedintrinsic point defects.

Finally, in a region ranging from an axially position from about 100 mmto about 125 mm there is an axially symmetric region of interstitialdominated material free of agglomerated defects surrounding a generallycylindrical core of vacancy dominated material. Within the vacancydominated material, there is an axially symmetric region which is freeof agglomerated defects surrounding a core containing agglomeratedvacancy defects.

Example 11 Cooling Rate and Position of V/I Boundary

A series of single crystal silicon ingots (150 mm and 200 mm nominaldiameter), were grown in accordance with the Czochralski method usingdifferent hot zone configurations, designed by means common in the art,which affected the residence time of the silicon at temperatures inexcess of about 1050° C. The pull rate profile for each ingot was variedalong the length of the ingot in an attempt to create a transition froma region of agglomerated vacancy point defects to a region ofagglomerated interstitial point defects.

Once grown, the ingots were cut longitudinally along the central axisrunning parallel to the direction of growth, and then further dividedinto sections which were each about 2 mm in thickness. Using the copperdecoration technique previously described, one set of such longitudinalsections was then heated and intentionally contaminated with copper, theheating conditions being appropriate for the dissolution of a highconcentration of copper interstitials. Following this heat treatment,the samples were then rapidly cooled, during which time the copperimpurities either outdiffused or precipitated at sites where oxideclusters or agglomerated interstitial defects where present. After astandard defect delineating etch, the samples were visually inspectedfor the presence of precipitated impurities; those regions which werefree of such precipitated impurities corresponded to regions which werefree of agglomerated interstitial defects.

Another set of the longitudinal sections was subjected to a series ofoxygen precipitation heat treatments in order to cause the nucleationand growth of new oxide clusters prior to carrier lifetime mapping.Contrast bands in lifetime mapping were utilized in order to determineand measure the shape of the instantaneous melt/solid interface atvarious axial positions in each ingot. Information on the shape of themelt/solid interface was then used, as discussed further below, toestimate the absolute value of, and the radial variation in, the averageaxial temperature gradient, G₀. This information was also used, inconjunction with the pull rate, to estimate the radial variation inv/G₀.

To more closely examine the effect growth conditions have on theresulting quality of a single crystal silicon ingot, several assumptionswere made which, based on experimental evidence available to-date, arebelieved to be justified. First, in order to simplify the treatment ofthermal history in terms of the time taken to cool to a temperature atwhich the agglomeration of interstitial defects occurs, it was assumedthat about 1050° C. is a reasonable approximation for the temperature atwhich the agglomeration of silicon self-interstitials occurs. Thistemperature appears to coincide with changes in agglomeratedinterstitial defect density observed during experiments in whichdifferent cooling rates were employed. Although, as noted above, whetheragglomeration occurs is also a factor of the concentration ofinterstitials, it is believed that agglomeration will not occur attemperatures above about 1050° C. because, given the range ofinterstitial concentrations typical for Czochralski-type growthprocesses, it is reasonable to assume that the system will not becomecritically supersaturated with interstitials above this temperature.Stated another way, for concentrations of interstitials which aretypical for Czochralski-type growth processes, it is reasonable toassume that the system will not become critically supersaturated, andtherefore an agglomeration event will not occur, above a temperature ofabout 1050° C.

The second assumption that was made to parameterize the effect of growthconditions on the quality of single crystal silicon is that thetemperature dependence of silicon self-interstitial diffusivity isnegligible. Stated another way, it is assumed that self-interstitialsdiffuse at the same rate at all temperatures between about 1400° C. andabout 1050° C. Understanding that about 1050° C. is considered areasonable approximation for the temperature of agglomeration, theessential point of this assumption is that the details of the coolingcurve from the melting point does not matter. The diffusion distancedepends only on the total time spent cooling from the melting point toabout 1050° C.

Using the axial temperature profile data for each hot zone design andthe actual pull rate profile for a particular ingot, the total coolingtime from about 1400° C. to about 1050° C. may be calculated. It shouldbe noted that the rate at which the temperature changes for each of thehot zones was reasonably uniform. This uniformity means that any errorin the selection of a temperature of nucleation for agglomeratedinterstitial defects, i.e. about 1050° C., will arguably lead only toscaled errors in the calculated cooling time.

In order to determine the radial extent of the vacancy dominated regionof the ingot (R_(vacancy)), or alternatively the width of the axiallysymmetric region, it was further assumed that the radius of the vacancydominated core, as determined by the lifetime map, is equivalent to thepoint at solidification where v/G₀=v/G₀ critical. Stated another way,the width of the axially symmetric region was generally assumed to bebased on the position of the V/I boundary after cooling to roomtemperature. This is pointed out because, as mentioned above, as theingot cools recombination of vacancies and silicon self-interstitialsmay occur. When recombination does occur, the actual position of the V/Iboundary shifts inwardly toward the central axis of the ingot. It isthis final position which is being referred to here.

To simplify the calculation of G₀, the average axial temperaturegradient in the crystal at the time of solidification, the melt/solidinterface shape was assumed to be the melting point isotherm. Thecrystal surface temperatures were calculated using finite elementmodeling (FEA) techniques and the details of the hot zone design. Theentire temperature field within the crystal, and therefore G₀, wasdeduced by solving Laplace's equation with the proper boundaryconditions, namely, the melting point along the melt/solid interface andthe FEA results for the surface temperature along the axis of thecrystal. The results obtained at various axial positions from one of theingots prepared and evaluated are presented in FIG. 27.

To estimate the effect that radial variations in G₀ have on the initialinterstitial concentration, a radial position R′, that is, a positionhalfway between the V/I boundary and the crystal surface, was assumed tobe the furthest point a silicon self-interstitial can be from a sink inthe ingot, whether that sink be in the vacancy dominated region or onthe crystal surface. By using the growth rate and the G₀ data for theabove ingot, the difference between the calculated v/G₀ at the positionR′ and v/G₀ at the V/I boundary (i.e., the critical v/G₀ value) providesan indication of the radial variation in the initial interstitialconcentration, as well as the effect this has on the ability for excessinterstitials to reach a sink on the crystal surface or in the vacancydominated region.

For this particular data set, it appears there is no systematicdependence of the quality of the crystal on the radial variation inv/G₀. As can be seen in FIG. 28, the axial dependence in the ingot isminimal in this sample. The growth conditions involved in this series ofexperiments represent a fairly narrow range in the radial variation ofG₀. As a result, this data set is too narrow to resolve a discernabledependence of the quality (i.e., the presence of absence of a band ofagglomerated intrinsic point defects) on the radial variation of G₀.

As noted, samples of each ingot prepared were evaluated at various axialpositions for the present or absence of agglomerated interstitialdefects. For each axial position examined, a correlation may be madebetween the quality of the sample and the width of the axially symmetricregion. Referring now to FIG. 29, a graph may be prepared which comparesthe quality of the given sample to the time the sample, at thatparticular axial position, was allowed to cool from solidification toabout 1050° C. As expected, this graph shows the width of the axiallysymmetric region (i.e., R_(crystal)−R_(vacancy)) has a strong dependenceon the cooling history of the sample within this particular temperaturerange. In order of the width of the axially symmetric region toincrease, the trend suggests that longer diffusion times, or slowercooling rates, are needed.

Based on the data present in this graph, a best fit line may becalculated which generally represents a transition in the quality of thesilicon from “good” (i.e., defect-free) to “bad” (i.e., containingdefects), as a function of the cooling time allowed for a given ingotdiameter within this particular temperature range. This generalrelationship between the width of the axially symmetric region and thecooling rate may be expressed in terms of the following equation:(R _(crystal) −R _(transition))² =D _(eff) *t _(1050° C.)wherein

R_(crystal) is the radius of the ingot,

R_(transition) is the radius of the axially symmetric region at an axialposition in the sample were a transition occurs in the interstitialdominated material from being defect-free to containing defects, or viceversa,

D_(eff) is a constant, about 9.3*10⁻⁴ cm²sec⁻¹, which represents theaverage time and temperature of interstitial diffusivity, and

t_(1050° C.) is the time required for the given axial position of thesample to cool from solidification to about 1050° C.

Referring again to FIG. 29, it can be seen that, for a given ingotdiameter, a cooling time may be estimated in order to obtain an axiallysymmetric region of a desired diameter. For example, for an ingot havinga diameter of about 150 mm, an axially symmetric region having a widthabout equal to the radius of the ingot may be obtained if, between thetemperature range of about 1410° C. and about 1050° C., this particularportion of the ingot is allowed to cool for about 10 to about 15 hours.Similarly, for an ingot having a diameter of about 200 mm, an axiallysymmetric region having a width about equal to the radius of the ingotmay be obtained if between this temperature range this particularportion of the ingot is allowed to cool for about 25 to about 35 hours.If this line is further extrapolated, cooling times of about 65 to about75 hours may be needed in order to obtain an axially symmetric regionhaving a width about equal to the radius of an ingot having a diameterof about 300 mm. It is to be noted in this regard that, as the diameterof the ingot increases, additional cooling time is required due to theincrease in distance that interstitials must diffuse in order to reachsinks at the ingot surface or the vacancy core.

Referring now to FIGS. 30, 31, 32 and 33, the effects of increasedcooling time for various ingots may be observed. Each of these figuresdepicts a portion of a ingot having a nominal diameter of 200 mm, withthe cooling time from the temperature of solidification to 1050° C.progressively increasing from FIG. 30 to FIG. 33.

Referring to FIG. 30, a portion of an ingot, ranging in axial positionfrom about 235 mm to about 350 mm from the shoulder, is shown. At anaxial position of about 255 mm, the width of the axially symmetricregion free of agglomerated interstitial defects is at a maximum, whichis about 45% of the radius of the ingot. Beyond this position, atransition occurs from a region which is free of such defects, to aregion in which such defects are present.

Referring now to FIG. 31, a portion of an ingot, ranging in axialposition from about 305 mm to about 460 mm from the shoulder, is shown.At an axial position of about 360 mm, the width of the axially symmetricregion free of agglomerated interstitial defects is at a maximum, whichis about 65% of the radius of the ingot. Beyond this position, defectformation begins.

Referring now to FIG. 32, a portion of an ingot, ranging in axialposition from about 140 mm to about 275 mm from the shoulder, is shown.At an axial position of about 210 mm, the width of the axially symmetricregion is about equal to the radius of the ingot; that is, a smallportion of the ingot within this range is free of agglomerated intrinsicpoint defects.

Referring now to FIG. 33, a portion of an ingot, ranging in axialposition from about 600 mm to about 730 mm from the shoulder, is shown.Over an axial position ranging from about 640 mm to about 665 mm, thewidth of the axially symmetric region is about equal to the radius ofthe ingot. In addition, the length of the ingot segment in which thewidth of the axially symmetric region is about equal to the radius ofthe ingot is greater than what is observed in connection with the ingotof FIG. 32.

When viewed in combination, therefore, FIGS. 30, 31, 32, and 33demonstrate the effect of cooling time to 1050° C. upon the width andthe length of the defect-free, axially symmetric region. In general, theregions containing agglomerated interstitial defects occurred as aresult of a continued decrease of the crystal pull rate leading to aninitial interstitial concentration which was too large to reduce for thecooling time of that portion of the crystal. A greater length of theaxially symmetric region means a larger range of pull rates (i.e.,initial interstitial concentration) are available for the growth of suchdefect-free material. Increasing the cooling time allows for initiallyhigher concentration of interstitials, as sufficient time for radialdiffusion may be achieved to suppress the concentration below thecritical concentration required for agglomeration of interstitialdefects. Stated in other words, for longer cooling times, somewhat lowerpull rates (and, therefore, higher initial interstitial concentrations)will still lead to the maximum axially symmetric region 60. Therefore,longer cooling times lead to an increase in the allowable pull ratevariation about the condition required for maximum axially symmetricregion diameter and ease the restrictions on process control. As aresult, the process for an axially symmetric region over large lengthsof the ingot becomes easier.

Referring again to FIG. 33, over an axial position ranging from about665 mm to greater than 730 mm from the shoulder of crystal, a region ofvacancy dominated material free of agglomerated defects is present inwhich the width of the region is equal to the radius of the ingot.

As can be seen from the above data, by means of controlling the coolingrate, the concentration of self-interstitials may be suppressed byallowing more time for interstitials to diffuse to regions where theymay be annihilated. As a result, the formation of agglomeratedinterstitial defects is prevented within significant portion of thesingle crystal silicon ingot.

In view of the above, it will be seen that the several objects of theinvention are achieved.

As various changes could be made in the above compositions and processeswithout departing from the scope of the invention, it is intended thatall matter contained in the above description be interpreted asillustrative and not in a limiting sense.

1. A process for heat-treating a Cz, single crystal silicon wafer toinfluence the precipitation behavior of oxygen in the wafer in asubsequent thermal processing step, the wafer comprising two major,generally parallel surfaces, one of which is a front surface of thewafer and the other of which is a back surface of the wafer, a centralplane between the front and back surfaces, a circumferential edgejoining the front and back surfaces, a central axis generallyperpendicular to the front and back surfaces, a radius extending fromthe central axis to the circumferential edge, a nominal diameter of 150mm, 200 mm, or greater than 200 mm, a front surface layer whichcomprises the region of the wafer between the front surface and adistance, D₁, of at least about 10 micrometers measured from the frontsurface and toward the central plane, a back surface layer whichcomprises a second region of the wafer between the back surface and adistance D₂, of at least about 10 micrometers as measured from the backsurface and toward the central plane, and a bulk layer which comprises athird region of the wafer between the front surface layer and the backsurface layer, wherein said wafer further comprises an axially symmetricregion which has radial width of at least about 30% the length of theradius of the wafer and has no detectible agglomerated intrinsic pointdefects at a detection limit of 10³ defects/cm³, the process comprisingthe steps of; subjecting the wafer to a heat-treatment in excess ofabout 1175° C. in a non-oxidizing atmosphere to form crystal latticevacancies in the front surface and bulk layers; and, controlling thecooling rate of the heat-treated wafer to produce a wafer having avacancy concentration profile in which a maximum concentration islocated between the central plane and the back surface layer, thevacancy concentration generally increasing from the front surface to theregion of maximum concentration and generally decreasing from the regionof maximum concentration to the back surface with the difference in theconcentration of vacancies in the front surface layer, the back surfacelayer, and the bulk layer being such that a thermal treatment at atemperature in excess of 750° C. is capable of forming in the wafer adenuded zone in the front surface layer and in the back surface layerand oxygen clusters or precipitates in the bulk zone, with theconcentration of the oxygen clusters or precipitates in the bulk layerbeing primarily dependant upon the concentration of vacancies.
 2. Theprocess of claim 1 wherein said heat-treatment to form crystal latticevacancies further comprises exposing the front surface to anon-nitriding atmosphere and the back surface to a nitriding atmosphere.3. The process of claim 1 wherein said heat-treatment to form crystallattice vacancies further comprises annealing the wafer with a secondsilicon wafer having a front surface that is generally parallel to aback surface, wherein the wafers are arranged such that the wafer'sfront surface is in contact with the second wafer's front surface in aface-to-face arrangement.
 4. The process of claim 2 wherein saidnitriding atmosphere is selected from a group consisting of nitrogen andammonia.
 5. The process of claim 4 wherein said nitriding atmosphere isnitrogen.
 6. The process of claim 5 wherein said nitriding atmospherefurther comprises oxygen.
 7. A process for producing a single crystalsilicon wafer having two major, generally parallel surfaces, one ofwhich is a front surface of the wafer and the other of which is a backsurface of the wafer, a central plane between the front and backsurfaces, a circumferential edge joining the front and back surfaces, acentral axis generally perpendicular to the front and back surfaces, aradius extending from the central axis to the circumferential edge, anominal diameter of 150 mm, 200 mm, or greater than 200 mm, a frontsurface layer which comprises the region of the wafer between the frontsurface and a distance, D₁, of at least about 10 micrometers measuredfrom the front surface and toward the central plane, a back surfacelayer which comprises a second region of the wafer between the backsurface and a distance D₂, of at least about 10 micrometers as measuredfrom the back surface and toward the central plane, and a bulk layerwhich comprises a third region of the wafer between the front surfacelayer and the back surface layer, the wafer being characterized in thatthe wafer has a non-uniform distribution of crystal lattice vacancies inwhich a maximum concentration is located between the central plane andthe back surface layer, the vacancy concentration generally increasingfrom the front surface to the region of maximum concentration andgenerally decreasing from the region of maximum concentration to theback surface, the wafer further comprising a first axially symmetricregion, which has no detectible agglomerated intrinsic point defects ata detection limit of 10³ defects/cm³, the process comprising the stepsof: growing a single crystal silicon ingot in which the ingot comprisesa central axis, a seed-cone, an end-cone and a constant diameter portionbetween the seed-cone and the end-cone having a circumferential edge, aradius extending from the central axis to the circumferential edge, anda nominal diameter of 150 mm, 200 mm, or greater than 200 mm;controlling (i) a growth velocity, v, (ii) an average axial temperaturegradient, G₀, during the growth of the constant diameter portion of thecrystal over the temperature range from solidification to a temperatureof no less than about 1325° C., and (iii) the cooling rate of thecrystal from the solidification temperature to about 1,050° C. to causethe formation of an axially symmetrical segment which has no detectibleagglomerated intrinsic point defects at a detection limit of 10³defects/cm³, wherein the axially symmetric region extends inwardly fromthe circumferential edge of the ingot, has a width as measured from thecircumferential edge radially toward the central axis of the ingot whichis at least about 30% the length of the radius of the ingot, and has alength as measured along the central axis of at least about 20% thelength of the constant diameter portion of the ingot; slicing a singlecrystal silicon wafer from the constant diameter portion of the ingot;subjecting the wafer to a heat-treatment in excess of about 1175° C. ina non-oxidizing atmosphere to form crystal lattice vacancies in thefront surface and bulk layers; and, controlling the cooling rate of theheat-treated wafer to produce a wafer having a vacancy concentrationprofile in which a maximum concentration is located between the centralplane and the back surface layer, the vacancy concentration generallyincreasing from the front surface to the region of maximum concentrationand generally decreasing from the region of maximum concentration to theback surface and the difference in the concentration of vacancies in thefront surface layer, the back surface layer, and the bulk layer beingsuch that a thermal treatment at a temperature in excess of 750° C. iscapable of forming in the wafer a denuded zone in the front surfacelayer and in the back surface layer and oxygen clusters or precipitatesin the bulk zone, with the concentration of the oxygen clusters orprecipitates in the bulk layer being primarily dependant upon theconcentration of vacancies.
 8. The process of claim 7 wherein saidheat-treatment to form crystal lattice vacancies further comprisesexposing the front surface to a non-nitriding atmosphere and the backsurface to a nitriding atmosphere.
 9. The process of claim 7 whereinsaid heat-treatment to form crystal lattice vacancies further comprisesannealing the wafer with a second silicon wafer having a front surfacethat is generally parallel to a back surface, wherein the wafers arearranged such that the wafer's front surface is in contact with thesecond wafer's front surface in a face-to-face arrangement.
 10. Theprocess of claim 8 wherein said nitriding atmosphere is selected from agroup consisting of nitrogen and ammonia.
 11. The process of claim 10wherein said nitriding atmosphere is nitrogen.
 12. The process of claim11 wherein said nitriding atmosphere further comprises oxygen.
 13. Aprocess for heat-treating a Cz, single crystal silicon wafer, toinfluence the precipitation behavior of oxygen in the wafer in asubsequent thermal processing step, the wafer comprising two major,generally parallel surfaces, one of which is a front surface of thewafer and the other of which is a back surface of the wafer, a centralplane between the front and back surfaces, a circumferential edgejoining the front and back surfaces, a central axis generallyperpendicular to the front and back surfaces, a radius extending fromthe central axis to the circumferential edge, a nominal diameter of 150mm, 200 mm, or greater than 200 mm, a front surface layer whichcomprises the region of the wafer between the front surface and adistance, D₁, of at least about 10 micrometers measured from the frontsurface and toward the central plane, a back surface layer whichcomprises a second region of the wafer between the back surface and adistance D₂, of at least about 10 micrometers as measured from the backsurface and toward the central plane, and a bulk layer which comprises athird region of the wafer between the front surface layer and the backsurface layer, wherein said wafer further comprises a first axiallysymmetric region in which vacancies are the predominant intrinsic pointdefect and which has no detectible agglomerated intrinsic point defectsat a detection limit of 10³ defects/cm³, wherein the first axiallysymmetric region comprises the central axis or has a width of at leastabout 15 mm the process comprising the steps of; subjecting the wafer toa heat-treatment in excess of about 1175° C. in a non-oxidizingatmosphere to form crystal lattice vacancies in the surface layer andthe bulk layer; and, controlling the cooling rate of the heat-treatedwafer to produce a wafer having a non-uniform concentration of vacanciesin which a maximum concentration is located between the central planeand the back surface layer, the vacancy concentration generallyincreasing from the front surface to the region of maximum concentrationand generally decreasing from the region of maximum concentration to theback surface such that, upon subjecting the wafer to an oxygenprecipitation heat treatment, a denuded zone is formed in the surfacelayer and oxygen clusters or precipitates are formed in the bulk layerwith the concentration of the oxygen clusters or precipitates in thebulk layer being primarily dependant upon the concentration ofvacancies.
 14. The process of claim 13 wherein said heat-treatment toform crystal lattice vacancies further comprises exposing the frontsurface to a non-nitriding atmosphere and the back surface to anitriding atmosphere.
 15. The process of claim 13 wherein saidheat-treatment to form crystal lattice vacancies further comprisesannealing the wafer with a second silicon wafer having a front surfacethat is generally parallel to a back surface, wherein the wafers arearranged such that the wafer's front surface is in contact with thesecond wafer's front surface in a face-to-face arrangement.
 16. Theprocess of claim 14 wherein said nitriding atmosphere is selected from agroup consisting of nitrogen and ammonia.
 17. The process of claim 16wherein said nitriding atmosphere is nitrogen.
 18. The process of claim17 wherein said nitriding atmosphere further comprises oxygen.